James L. Smialek

Glenn Research Center, Cleveland, Ohio

Anita Garg

The University of Toledo, Toledo, Ohio

Microstructure and Oxidation of a MAX

Phase/Superalloy Hybrid Interface

NASA/TM—2014-216679

July 2014
1$6$67,3URJUDPLQ3UR¿OH

6LQFHLWVIRXQGLQJ1$6$KDVEHHQGHGLFDWHGWRWKH

DGYDQFHPHQWRIDHURQDXWLFVDQGVSDFHVFLHQFH7KH

1$6$6FLHQWL¿FDQG7HFKQLFDO,QIRUPDWLRQ67,

SURJUDPSOD\VDNH\SDUWLQKHOSLQJ1$6$PDLQWDLQ

WKLVLPSRUWDQWUROH

7KH1$6$67,3URJUDPRSHUDWHVXQGHUWKHDXVSLFHV

RIWKH$JHQF\&KLHI,QIRUPDWLRQ2I¿FHU,WFROOHFWV

RUJDQL]HVSURYLGHVIRUDUFKLYLQJDQGGLVVHPLQDWHV

1$6$¶V67,7KH1$6$67,SURJUDPSURYLGHVDFFHVV

WRWKH1$6$$HURQDXWLFVDQG6SDFH'DWDEDVHDQG

LWVSXEOLFLQWHUIDFHWKH1$6$7HFKQLFDO5HSRUWV

6HUYHUWKXVSURYLGLQJRQHRIWKHODUJHVWFROOHFWLRQV

RIDHURQDXWLFDODQGVSDFHVFLHQFH67,LQWKHZRUOG

5HVXOWVDUHSXEOLVKHGLQERWKQRQ1$6$FKDQQHOV

DQGE\1$6$LQWKH1$6$67,5HSRUW6HULHVZKLFK

LQFOXGHVWKHIROORZLQJUHSRUWW\SHV

 7(&+1,&$/38%/,&$7,215HSRUWVRI

FRPSOHWHGUHVHDUFKRUDPDMRUVLJQL¿FDQWSKDVH

RIUHVHDUFKWKDWSUHVHQWWKHUHVXOWVRI1$6$

SURJUDPVDQGLQFOXGHH[WHQVLYHGDWDRUWKHRUHWLFDO

DQDO\VLV,QFOXGHVFRPSLODWLRQVRIVLJQL¿FDQW

VFLHQWL¿FDQGWHFKQLFDOGDWDDQGLQIRUPDWLRQ

GHHPHGWREHRIFRQWLQXLQJUHIHUHQFHYDOXH

1$6$FRXQWHUSDUWRISHHUUHYLHZHGIRUPDO

SURIHVVLRQDOSDSHUVEXWKDVOHVVVWULQJHQW

OLPLWDWLRQVRQPDQXVFULSWOHQJWKDQGH[WHQWRI

JUDSKLFSUHVHQWDWLRQV

 7(&+1,&$/0(025$1'806FLHQWL¿F

DQGWHFKQLFDO¿QGLQJVWKDWDUHSUHOLPLQDU\RU

RIVSHFLDOL]HGLQWHUHVWHJTXLFNUHOHDVH

UHSRUWVZRUNLQJSDSHUVDQGELEOLRJUDSKLHVWKDW

FRQWDLQPLQLPDODQQRWDWLRQ'RHVQRWFRQWDLQ

H[WHQVLYHDQDO\VLV

 &2175$&7255(32576FLHQWL¿FDQG

WHFKQLFDO¿QGLQJVE\1$6$VSRQVRUHG

FRQWUDFWRUVDQGJUDQWHHV

 &21)(5(1&(38%/,&$7,21&ROOHFWHG

SDSHUVIURPVFLHQWL¿FDQGWHFKQLFDO

FRQIHUHQFHVV\PSRVLDVHPLQDUVRURWKHU

PHHWLQJVVSRQVRUHGRUFRVSRQVRUHGE\1$6$

 63(&,$/38%/,&$7,216FLHQWL¿F

WHFKQLFDORUKLVWRULFDOLQIRUPDWLRQIURP

1$6$SURJUDPVSURMHFWVDQGPLVVLRQVRIWHQ

FRQFHUQHGZLWKVXEMHFWVKDYLQJVXEVWDQWLDO

SXEOLFLQWHUHVW

 7(&+1,&$/75$16/$7,21(QJOLVK

ODQJXDJHWUDQVODWLRQVRIIRUHLJQVFLHQWL¿FDQG

WHFKQLFDOPDWHULDOSHUWLQHQWWR1$6$¶VPLVVLRQ

6SHFLDOL]HGVHUYLFHVDOVRLQFOXGHFUHDWLQJFXVWRP

WKHVDXULEXLOGLQJFXVWRPL]HGGDWDEDVHVRUJDQL]LQJ

DQGSXEOLVKLQJUHVHDUFKUHVXOWV

)RUPRUHLQIRUPDWLRQDERXWWKH1$6$67,

SURJUDPVHHWKHIROORZLQJ

 $FFHVVWKH1$6$67,SURJUDPKRPHSDJHDW

http://www.sti.nasa.gov

 (PDLO\RXUTXHVWLRQWRhelp@sti.nasa.gov

 )D[\RXUTXHVWLRQWRWKH1$6$67,

,QIRUPDWLRQ'HVNDW±±

 3KRQHWKH1$6$67,,QIRUPDWLRQ'HVNDW

 ±±

 :ULWHWR

67,,QIRUPDWLRQ'HVN

 1$6$&HQWHUIRU$HUR6SDFH,QIRUPDWLRQ

6WDQGDUG'ULYH

+DQRYHU0'±
James L. Smialek

Glenn Research Center, Cleveland, Ohio

Anita Garg

The University of Toledo, Toledo, Ohio

Microstructure and Oxidation of a MAX

Phase/Superalloy Hybrid Interface

NASA/TM—2014-216679

July 2014

1DWLRQDO$HURQDXWLFVDQG

6SDFH$GPLQLVWUDWLRQ

*OHQQ5HVHDUFK&HQWHU

&OHYHODQG2KLR
Acknowledgments

7KHDXWKRUVDUHJUDWHIXOWR'RQDOG+XPSKUH\IRUSURGXFLQJWKHKRWSUHVVHGVDPSOHV

-R\%XHKOHUIRUPHWDOORJUDSK\DQG7LP*DEEDQG&KDQWDO6XGEUDFNIRUKHOSIXOGLVFXVVLRQV

$YDLODEOHIURP

1$6$&HQWHUIRU$HURVSDFH,QIRUPDWLRQ

6WDQGDUG'ULYH

+DQRYHU0'±

1DWLRQDO7HFKQLFDO,QIRUPDWLRQ6HUYLFH

6KDZQHH5RDG

$OH[DQGULD9$

$YDLODEOHHOHFWURQLFDOO\DWKWWSZZZVWLQDVDJRY

7UDGHQDPHVDQGWUDGHPDUNVDUHXVHGLQWKLVUHSRUWIRULGHQWL¿FDWLRQ

RQO\7KHLUXVDJHGRHVQRWFRQVWLWXWHDQRI¿FLDOHQGRUVHPHQW

HLWKHUH[SUHVVHGRULPSOLHGE\WKH1DWLRQDO$HURQDXWLFVDQG

6SDFH$GPLQLVWUDWLRQ

7KLVZRUNZDVVSRQVRUHGE\WKH)XQGDPHQWDO$HURQDXWLFV3URJUDP

DWWKH1$6$*OHQQ5HVHDUFK&HQWHU

Level of Review7KLVPDWHULDOKDVEHHQWHFKQLFDOO\UHYLHZHGE\WHFKQLFDOPDQDJHPHQW
NASA/TM—2014-216679 1 

 Microstructure and Oxidation of a MAX 

Phase/Superalloy Hybrid Interface 

James L. Smialek 

National Aeronautics and Space Administration 

Glenn Research Center 

Cleveland, Ohio 44135 

Anita Garg 

The University of Toledo 

Toledo, Ohio 43606 

Abstract 

Corrosion resistant, strain tolerant MAX phase coatings are of interest for turbine applications. Thin 

Cr2AlC MAX phase wafers were vacuum diffusion bonded to an advanced turbine disk alloy, LSHR, at 

1100 qC. The interface, examined by optical and scanning electron microscopy, revealed a primary 

diffusion zone consisting of ~ 10 μm of E-Ni(Co)Al, decorated with various NiCoCrAl, MC and M3B2

precipitates. On the Cr2AlC side, an additional ~40 μm Al-depletion zone of Cr7C3 formed in an 

interconnected network with the E-Ni(Co)Al. Oxidation of an exposed edge at 800 qC for 100 h produced 

a fine-grained lenticular alumina scale over Cr2AlC and E-Ni(Co)Al, with coarser chromia granules over 

the Cr7C3 regions. Subsequent growth of the diffusion layers was only ~ 5 μm in total. A residual stress of 

~ 500 MPa was estimated for the MAX phase layer, but no interfacial damage was observed. Subsequent 

tests for 1000 h reveal similar results. 

Introduction 

M-A-X compounds are carbide or nitride ceramics well known for their unique behavior as 

‘deformable’ and ‘strain tolerant’ ceramics. This property arises from the special hexagonal crystal 

structure that contains weakly bonded (0003) hkil crystallographic planes.(Barsoum & El-raghy, 2001) 

(Eklund, Beckers, Jansson, Högberg, & Hultman, 2010). These can deform by sliding and kinking rather 

than by the macroscopic cracking and cleavage typical of most ceramics. The chemistry is generically 

described as M = group III-VI transition metals (Sc, Ti, V, Cr); A = group IIIA –VA main elements (Al, 

Si, Ge, P); and X = C, N. The stoichiometry is commonly M2AX, M3AX2, or M4AX3. The structure is 

distinguished by planes of A-group elements every 3rd, 4th, or 5th layer, respectively, in a global Mn+1Xn

lattice. While many of the M-X ceramic properties are maintained (high melting, high modulus, high 

strength), the relatively weak M-A and A-X bonds allow for easy intercalated planar delamination that 

results in high strain tolerance. This produces unique aspects such as high machinability, fracture 

toughness, thermal fatigue resistance, and coefficient of thermal expansion for an otherwise ‘ceramiclike’ material. 

While there have been over 60 MAX phases identified, those containing Al have the potential of 

excellent oxidation resistance owing to the formation of Al2O3 scales (Tallman, Anasori, & Barsoum, 

2013). These include Ti2AlN, Ti3AlC2, Ti2AlC, and Cr2AlC, the latter two exhibiting among the highest 

oxidation resistance (Wubian Tian, Wang, Kan, & Zhang, 2008) (Z. J. Lin, Li, Wang, & Zhou, 2007) 

(D.B. Lee, Nguyen, Han, & Park, 2007). In general these studies have found alumina scale growth rates in 

the range of those formed on oxidation resistant NiAl or FeCrAl alloys, with good oxidation resistance up 
NASA/TM—2014-216679 2 

to 1300 qC for Ti3AlC2 and Ti2AlC and 1200 qC for Cr2AlC (Tallman et al., 2013). Ti2AlC has been 

characterized as having good cyclic oxidation resistance because of its good thermal expansion match to 

that of the alumina scale. The higher CTE of Cr2AlC leads to an increased spallation tendency similar to 

that observed for metals, but at much higher temperatures. 

Furthermore, studies document good corrosion resistance of Ti3AlC2 and Cr2AlC in SO2

environments (Dong Bok Lee & Park, 2011) (Dong Bok Lee, Nguyen, & Park, 2011). Also, hot corrosion 

resistance was demonstrated in molten Na2SO4 salt for Cr2AlC (Z. Lin, Zhou, Li, & Wang, 2006a), and 

for Ti3AlC2, but only if preoxidized (Z. Lin, Zhou, Li, & Wang, 2006b). Thus the production of MAX 

phases as protective coatings for stainless steel, Ni-base superalloy M38G, Ti6242, and TiAl substrates 

has been demonstrated, primarily via magnetron sputtering (Eklund et al., 2010) (Walter, Sigumonrong, 

El-Raghy, & Schneider, 2006)( (Gulbiski, 2004)(Hajas et al., 2011) (Wang, Flores Renteria, et al., 

2010)(Wang, Mykhaylonka, et al., 2010) (Li, Li, Xiang, Lu, & Zhou, 2011). 

Emerging problem areas of environmental degradation in turbines are gas phase embrittlement, 

oxidation, and low temperature hot corrosion (LTHC) of advanced disk alloys used in the high pressure 

turbine stage. (Woodford, 2006) (Birbilis & Buchheit, 2008)(Sudbrack et al., 2012) (Encinas-Oropesa et 

al., 2008)(Sumner, Encinas-Oropesa, Simms, & Oakey, 2011) (Karabela et al., 2011)(Sato, Chiu, & Reed, 

2011)(Moverare & Johansson, 2010) (Cruchley, Evans, Taylor, Hardy, & Stekovic, 2013). This is an 

especially sensitive application in that the surface of the disk and blade attachment points represent the 

most highly stressed areas of a turbine. Accordingly, crack tip oxidation has experienced in-depth 

attention (Kitaguchi et al., 2013)(Karabela, Zhao, Lin, Tong, & Hardy, 2013). Typical metallic aluminide 

and NiCrAlY coatings, engineered for higher temperature blade exposures and Type II hot corrosion 

resistance, actually lead to fatigue debits because of CTE mismatch stresses, lower strength, and tendency 

for brittle behavior in a disk environment. This warrants modified approaches, such as fully ductile Ni-Cr 

coatings or fatigue resistant, refractory strengthened NiCrAl-Ta,W “EQ” coating alloys (Mercer, 

Kawagishi, Tomimatsu, Hovis, & Pollock, 2011). 

We therefore propose that Cr2AlC may be worth considering as a potential coating in this application. 

It has three desirable attributes: strain tolerance due to microlaminate kinking, a relatively high CTE 

(13u10-6/qC), and good Type I hot corrosion resistance. Our current studies indicate good Type II low 

temperature corrosion behavior as well. However little information is available regarding cyclic thermal 

stability of this coating with a superalloy. Thus the purpose of the present paper is to examine the 

compatibility of a hot pressed Cr2AlC-LSHR disk alloy hybrid after repeated cycling to 800 qC. 

Interfacial mechanical, oxidative, and diffusional stabilities are the primary focus points. 

Materials and Procedure 

The alloy portion of the couple was LSHR (Low J’ Solvus temperature, High Refractory) developed 

by NASA for disk applications (Timothy P Gabb, Gayda, Telesman, & Kantzos, 2005). The AF-LSHR 

composition is:

 Ni-20.4Co-12.3Cr-4.3W-3.4Ti-3.4Al-2.7Mo-1.5Ta-1.5Nb-0.05Zr-.03B-.04C. 

Production scale powder metallurgy disks were produced by argon atomized, consolidated, hot 

compacted powders, followed by extrusion and isothermal forging. Specimen blanks were supersolvus 

heat treat at 1171 qC for 2 h, cooled at 72 qC/min., then aged at 855 qC/4 h and 775 qC/8 h. Phase

constituents were primarily J-Ni(Co,Cr,Mo,W) solid solution, strengthened by dense Jc-Ni(Ti,Ta,Nb)3Al 

precipitates, and dispersed (Ta,Nb,Ti)C carbide, coarse (W,Mo,Cr)3B2 boride, and fine grain boundary 

(W,Cr)23C6 carbide particles. (T P Gabb & Miller, 2012).
NASA/TM—2014-216679 3 

The starting Cr2AlC MAX phase ingot, approximately 2- by 2- by 12-cm, was obtained from 

Sandvik/Kanthal. As-received density was ~67% based on an assumed theoretical density of 5.22 g/cm3

. 

Thus an effort was made to improve density by hot pressing in vacuum (77 Pa or 10-6 torr) using graphite 

dies. It was found that hot pressing at 1300 qC for 2 h using 35 MPa pressure resulted in 97% density. The 

phase constituency was estimated by XRD Reitveld analyses to be ~92.3% Cr2AlC, 3.9% Al2O3, and 

3.8% Cr7C3 which did not change appreciably in bulk with 1300 qC hot pressing, 1100 qC diffusion 

bonding, or 800 qC oxidation. 

Coupons ~ 6- by 12-mm were sectioned from the hot pressed Cr2AlC and LSHR using a diamond 

wafer saw. Thickness was ~ 0.3 and 1.8 mm, respectively. These were polished to 2400 grit SiC emery 

finish and vacuum hot pressed at 1100 qC for 4 h under 85 MPa pressure (DC2). Hot pressing caused 

substantial deformation of the LSHR alloy, to the extent that the slightly undersized Cr2AlC layer was 

fully impressed into the metal. For this reason a second couple (DC3) was prepared for a subsequent 

1000 h oxidation test, using 1.5 mm Cr2AlC and only 50 MPa pressure. Only minimal deformation of the 

LSHR alloy was noted. Grafoil mold release was burned off at 800 qC for 1 h, and one end of the couple 

was sectioned for as-diffusion bonded metallography. 

The cut edge and major surfaces were polished to 4000 grit emery and ultrasonically cleaned in 

ethanol. The couple was oxidized at 800 qC in air for 100 h in a Thermolyne resistance muffle furnace, 

with intermittent removal for weight change and optical photography. The materials were analyzed by 

conventional metallographic, X-ray diffraction (XRD), and scanning electron microsopy (SEM) 

techniques before and after oxidation. SEM samples were coated with a conductive carbon or Au-Pd 

coating. Because of fine particle sizes and carbon coating, more quantitative assessments of carbide 

phases and thin oxide nodules were not attempted. Vacuum infiltrated epoxy mounted and polished crosssections were prepared before and after oxidation. An additional feature of this study was the examination 

of surface scales formed on the edge, which had been pre-polished to 4000 grit emery, to show 

microstructural evolution of the scales across the interfacial bond. 

Results 

The weight change behavior of the diffusion bonded hybrid couples are shown in Figure 1(a). Very 

little oxidation occurred, gaining only 0.09 mg/cm2

 after 100 h. This would be equivalent to a low 

parabolic rate constant of 7.23u10–5 mg2

/cm4/h for the couple, although cubic kinetics may ultimately be a 

more accurate description of layer growth rates (Tallman et al., 2013). However, both the Cr2AlC and 

LSHR alloy contribute to this value, so it is impossible to precisely characterize the weight change 

contribution of either. The weight change of the second couple was oxidized in a similar matter and 

achieved only 0.17 mg/cm2

 after 1000 h. There were no signs of delamination or deterioration of the bond 

interface. The literature suggests weight changes on the order of 0.061 and 0.20 for Cr2AlC oxidized for 

100 and 1000 h, respectively,(Z. J. Lin et al., 2007) and ~0.31 and 0.57 mg/cm2

 for LSHR oxidized for 

100 and 1000 h, respectively. (T. P. Gabb, Sudbrack, Draper, MacKay, & Telesman, 2014). Since weight 

gain is not a focus of this study, the kinetics are not discussed in more detail. 

XRD analyses of the oxidized surfaces is presented in Table 1. The Cr2AlC side formed primarily 

D-Al2O3, with an increase in the Cr7C3 phase. Also, another Cr3C2 carbide depletion phase was identified. 

The LSHR side formed primarily an D-Cr2O3 scale with some (Ti,Ta,Cr)O6 tetragonal rutile-type oxide. 
NASA/TM—2014-216679 4 

TABLE 1.—XRD SURFACE ANALYSIS OF HYBRID COUPLE OXIDIZED 100 h AT 800 qC 

MAX phase 

Chemical formula Compound name Crystal system Ref. code Relative intensity 

Cr2AlC MAX phase Hexagonal 04-007-2697 Strong

Al2O3 D-alumina Rhombohedral 01-089-7716 Weak 

Cr7C3 chrome carbide Orthorhombic 01-089-5902 Strong

Cr3C2 chrome carbide Orthorhombic 04-004-4541 Weak 

LSHR 

Chemical formula Compound name Crystal system Ref. code Relative intensity 

Ni(Nb) J-Ni solid solution Cubic 04-003-2243 Strong

(Ti,Ta,Cr)O6 rutile Tetragonal 01-076-7954 Weak 

(Cr,Ti)2O3 D-chromia Rhombohedral 01-082-0211 Medium 
NASA/TM—2014-216679 5 

The bonded interface of the diffusion couples are presented below as comparisons of polished crosssections of the hot pressed structure (HP) and the heat treated structure (OX), and the oxidized edge 

(EDGE). The cut and oxidized edge is shown in the optical micrographs of Figure 2. The full thickness of 

the Cr2AlC layer can be seen to be on the order of 250 μm, with an overall diffusion zone of ~ 50 μm. 

There is no appreciable change from 20 to 100 h of oxidation. Nor are any edge cracks or delamination 

visible. This indicates the relatively stable nature and low CTE mismatch stress of this couple. Optical 

micrographs are presented in Figure 3 for the mounted and polished samples, before (HP) and after (OX) 

the thermal treatment. A complex, multiphase interdiffusion zone can be seen, without major change due 

to the thermal exposure. Detailed SEM and energy dispersive spectroscopy (EDS) characterization is thus 

presented below for all three cases. 
NASA/TM—2014-216679 6 

Away from the interface, LSHR and Cr2AlC structures are characterized as unaffected endpoints. 

LSHR presented in Figure 4 shows densely populated, fine Jc-Ni3Al precipitates (#5,2) in both HP and 

OX polished samples. They are only slightly distinguished in baskscattered electron (BSE) images (4) 

because of the atomic weight similarity to the J-Ni(ss) matrix, and not appreciably altered by the 800 qC, 

100 h oxidation treatment. In contrast, sparsely distributed submicron (Ta,Nb,Ti)C carbide particles 

(#4,1) are easily distinguished because of their high refractory metal content. The corresponding LSHR 

region of the oxidized edge exhibits a profusion of Al,Cr,Ti,Ni-rich oxide nodules with faceted and 

needle-like surface crystals, Figure 5, with no obvious correlation with substrate microstructure. The 

various features characterized by EDS obtained at 15 kV are summarized in Table 2 and catalogued by 

feature number and exposure. Here EDS relative peak intensities are indicated in order of decreasing 

level, with ! indicating a difference of more than 15% and !! indicating a difference of more than 50%. 

Representative spectra will be presented later. Thus the unaffected LSHR alloy exhibits a Ni-rich matrix 

with low levels of all the alloying elements (#5), while the various carbides are rich in Ti, Cr, or 

refractory metals (#4). The LSHR scale features were rich in Al, Cr, Ti, Ni. Except for the Al-rich scale, 

this was consistent with XRD analyses that identified Cr(Ti)2O3 corundum patterns. 
TABLE 2.—SEM/EDS FEATURE ANALYSIS AND EXPOSURE EFFECTS (COLOR CODING BY SIMILAR ITEMS). FEATURES ARRANGED 

. Point Polished cross-sections: 

HP P025 As-Hot Pressed 

Point Polished cross-sections: 

OX P026 After Oxidation 

Point Polished 

Oxidized 

LSHR 5 matrix Ni>>(Co,Cr,Al,W,C,Mo,Ti) J/Jc 1 g.b. light particle Ta>(Ti,Nb)>C>(Cr,Ni,Co) MC 0A facets, needles O>(Al,Cr)>(Ti,NLSHR 4 grey particle Ta>(Ni,Nb,Ti)>C>(Al,Cr,Co) MC 2 g.b. precipitate Ni>>(Al,Ti,W,Co,Cr) JcLSHRDZ 9 grey particle Ta>(Nb,Ti)>>C MC 6 grey particle Ta>(Nb,Ti)>C>(Cr,Ni,Co) MC A nodule needles O>Ti>(Al,)Cr>>LSHRDZ 7 needle Ni>>(Al,Cr)>(W,Mo,Ti) TCP B flat base Al>(Cr,Ni)>(CoLSHRDZ 8 matrix Ni>>Cr>(Al,Co)> Ti J/Jc C nodule needles O>(Ti,Al,Cr)>(NDZLSHR 6 white particle W>>(Mo,Cr)>(Ti,C) M3B2 5 D.Z matrix (Ni,Al)>>>(Co,Cr,Ti) E-NiAl G nodule O>Cr>Ti>(Ta,ADZLSHR 7 grey particle Ta>(Ti,Nb)>>(Ni,Al,C) MC 3 white particle W>(Mo,Cr)>C M3B2 H faceted nodule O>(Al,Cr,Ti)>>DZLSHR 8 underlayer Ni>Al>>(Cr,Co,Ti) E-NiAl 4 grey particle Ni>Al>(Cr,Co)>(W,Mo,Ti) Jc D flat base Al>O>(Cr,Ni,CDZ1Cr2AlC 13 dark particle Al>>O>(Ni,Ti,Cr,Co) Al2O3 9 dark bond line particle Al>>O>(Ti,Ta,Nb,Cr,Co) Al2O3 E flat layer Al>>(O,Cr,Ni)>DZ1Cr2AlC 12 light phase (Ta,Ti)>(Al,Nb)>(Ni,C) MC 11 light phase (Ta,Ti)> (Al,Nb)>(Ni,C)>Cr,Co MC F smooth nodule (Al,O)>Cr>Ti>(DZ1Cr2AlC 11 underlayer Ni>Al (Cr,Co) J/Jc 10 grey particle Ti>W>Nb>Mo>(C,Al)>>(Ni,Co,Cr) MC I sea urchin Al>O>Ti>Cr>(NDZ2Cr2AlC 10 amorphic islands Cr>>(W,Mo)>C Cr7C3 14 underlayer Ni>>Al >>(Ti,Cr,Co)>(W,Mo) J/Jc J amorphic fingers Al>O>Cr>(Ni,CDZ2Cr2AlC 15 between carbide network (Ni,Al)>>(Cr,Co) E-NiAl K granular islands (Cr,O)>Al>(Ni,DZ2Cr2AlC 12 amorphic islands Cr> (W,Mo)>C Cr7C3 M amorphic fingers Al>O>Cr>(Ni,CDZ2Cr2AlC 13 dark particle Al>>O Al2O3 L granular islands (Cr,O)>>Al DZ2Cr2AlC 16 outer D.Z. matrix (Ni,Al)>>(Co,Cr) E-NiAl N smooth transition Al>O>>Cr>(NiDZ2Cr2AlC 17 amorphic islands Cr>C>Al Cr7C3Cr2AlC 1 light phase Cr>>C Cr7C3 18 light phase Cr>>C Cr7C3 Q granular Cr>O Cr2AlC 2 dark particle Al>>O Al2O3 19 dark particle Al>>O Al2O3 O dark particle Al>>O Cr2AlC 3 matrix Cr=Al>>C Cr2AlC 20 matrix (Cr,Al)>>C Cr2AlC P smooth matrix Al>(Cr,O) 

NASA/TM—2014-216679 7 
NASA/TM—2014-216679 8 

For the HP sample away from the interface at the other extreme, the Cr2AlC layer can be seen in 

Figure 6 as the three-phase material previously indicated. The matrix phase here is Cr2AlC (#3), the light 

phase particles are Cr7C3 (#1), and the dark particles are Al2O3 (#2), consistent with the associated EDS 

peak intensities catalogued in Table 2. The matrix phase on the sample edge was oxidized to an Al(Cr)2O3

lenticular scale structure, Figure 7(a), similar to that reported previously for 900 qC oxidation. (D.B. Lee 

et al., 2007). However, a coarser Cr2O3 granular scale was observed in local areas probably corresponding 

to and growing on the Cr7C3 impurity particles, Figure 7(b). The more open structure here may be a result 

of CO/CO2 gas evolution occurring during oxidation. 

The as-hot pressed polished cross-section (HP) in Figure 8(a) shows ~260 μm thick Cr2AlC layer 

remaining after removing 40 μm during surface preparation and a primary diffusion layer of ~47 μm. 

More subtle gradients beyond this primary zone can also be distinguished. The as-hot pressed crosssection is compared to that after thermal exposure (8c, OX) and to the oxidized edge (8b, EDGE). The 

primary NiAl + Cr7C3 diffusion zone appears to have grown by an additional 4 to 5 μm. Thicknesses of 

the various diffusion zones characterized below are summarized in Table 3. 

TABLE 3.—DIFFUSION LAYER THICKNESSES (μm) 

LSHR 

carbide growth 

LSHR 

Ni-Al 

Cr2AlC 

Ni-Al 

Cr2AlC 

carbide growth 

'DZ 

Polished hot pressed HP 81.0 5.8 5.4 39.3 ---- 

Polished oxidized OX 77.7 7.2 6.8 42.6 +4.7 

As-oxidized EDGE (22.8) 6.4 6.5 41.1 +2.4 
NASA/TM—2014-216679 9 
NASA/TM—2014-216679 10 

Interfacial details are further compared in Figures 9 to 11 with EDS descriptions of labeled specific 

features listed in Table 2. On the LSHR alloy side (left) for the as-hot pressed interface (HP), Ta(Ti,Nb) 

rich carbides (#9,7) are observed in both the Ni-rich alloy matrix (#5) and in the NiAl primary diffusion 

zone (#8), Figure 9. Larger, bright W-rich borides (#6) are also in this region. The original bond line 

(B.L.) is indicated by the dashed yellow line, decorated by the fine dark Al2O3 precipitates (#13). Moving 

toward the Cr2AlC layer, light (Ta,Ti)C carbides (#12) are again seen in an NiAl matrix (#11). Moving 

further right, a high volume percent of a light amorphic Cr-rich phase is present, probably a diffusional 

growth of the Cr7C3 impurity phase. This phase is well known to be a sublayer diffusion zone that grows 

under the alumina scale as Al is extracted from Cr2AlC during oxidation. (D.B. Lee et al., 2007)(Z. J. Lin 

et al., 2007). Also, it has been predicted thermodynamically and shown experimentally that NiAl will 

deplete Al from Cr2AlC to form Cr7C3 (Hajas, et al. 2010). Given that the high content of Al in Cr2AlC is 

likely to diffuse into the Ni-based LSHR alloy, Cr7C3 might also be expected as a depletion zone in an 

LSHR- Cr2AlC couple. 
NASA/TM—2014-216679 11 
NASA/TM—2014-216679 12 

Figure 10 shows the interfacial details after oxidation in a mounted and polished cross section (OX). 

Features similar to those before oxidation are identified: Ta(Ti,Nb) MC carbides in LSHR (#6,11), Ni-Al 

primary diffusion zone (#5), bright W-rich borides (#3), fine Al2O3 precipitates (#9), and amorphic Cr7C3

colonies (#12). Some additional features are now apparent. A Cr(W,Mo,Ti)- rich acicular phase (#7) is 

present in the LSHR side, probably V. And a weakly Widmanstätten Ni-Al precipitate (#4,14) appeared in 

the Ni-Al primary diffusion zone (#5,15). Semi-quantitative analysis of the primary diffusion zone 

yielded Ni-13Co-38Al, consistent with hypostoichiometric E-NiAl. A grey carbide particle near the 

LSHR interface was estimated as 37Ti-17Ta-26Nb (carbon excluded from analysis). The fine d0.5 μm 

precipitates in the NiAl matrix varied considerably, with typical estimates near 30Ni-20Co-20Al-20Cr3Ti, but not immediately suggestive of any common alloy phase. 

Scale features of this interface area in the exposed edge are presented in Figure 11. Although a planar 

boundary can be discerned, there is no sign of bond line degradation. The fine particles at the bond line 

(B.L.) and many of the dark Al2O3 precipitates appear to have been masked by the scale features, while 

other structural variations and gradients remain. In general, flat Al-rich oxide (#D,E) formed over the 

NiAl diffusion zone with numerous nodule-type structures. On the LSHR side, the nodules (#G,H) were 

more often enriched in Cr, Ti, while on the Cr2AlC side, they were primarily Al-rich (#F,I). Figure 12 

presents a high magnification image of the Al-oxide needle colony (#I) present in Figure 11. While some 

rounding is apparent at the needle tips, spherical balls associated with a VLS process were not present at 

the ends. Figure 13(a) presents the interlocking structures of the Ni-Al matrix (#16) and large area 

fraction of the Cr-carbide (#17) within the diffusion zone produced by hot pressing. Figures 13(b) and (c) 

are from the oxidized edge, showing smooth Al-rich (#J,M,N) and granular Cr-rich oxide (#K,L), 

respectively, formed on these phases near the middle (13b) and Cr2AlC region (13c,d) of the diffusion 

zone. 
NASA/TM—2014-216679 13 
NASA/TM—2014-216679 14 

Representative EDS spectra are compared for the HP, OX, and EDGE samples in Figures 14 to 16 for 

the unaffected LSHR layer, the Ni-Al diffusion zone matrix phase, and the Cr-carbide phase near the 

Cr2AlC layer, respectively. Strong Al, Cr, and Ti enrichments are seen in the oxidized LSHR surface 

(Figure 14(c)), as compared to Ni-rich polished cross-sections (Figures 14(a) and (b)). Little change 

occurred in the Ni-Al diffusion zone after oxidation (Figures 15(a) and (b)), but the scale formed on this 

phase was very Al-rich, with only some Cr. Finally, the Cr-rich carbide phase in the diffusion zone 

changed little and oxidized to the Cr(Al)-rich oxide, Figure 16, interestingly without any Al evident in the 

starting phase. 

 
NASA/TM—2014-216679 15 
NASA/TM—2014-216679 16 

Discussion 

Thus a dense wafer of Cr2AlC has been successfully bonded to an advanced disk alloy, LSHR, at 

1100 qC. A noticeable amount of interdiffusion took place, producing a distinct interface layer of an NiAl rich zone, presumably E-NiAl. This zone appeared to straddle the original interface, as demarked by 

alumina stringers. Little growth took place after 100 h exposure to 800 qC ambient air. No evidence of 

interfacial damage or cracking was apparent. Various modifications of the interface zone were manifested 

as (Ta,Ti,Nb)C carbide particles and a profusion of (Ni,Co,Al,Cr)-rich Widmanstätten precipitates. 

Further in toward the Cr2AlC portion, a multiphase zone comprised of a large component of Cr-carbide 

has formed, analogous to the Cr7C3 depletion phase widely observed under alumina scales. The LSHR 

portion exhibited a dispersion of a V TCP needles, with a broad, fading zone of enlarged MC carbides. 

The characterization of the oxidized exposed edge reveals the most direct effect of the NiAl and 

Cr7C3 zones on oxide scale formation. That is, a relatively flat, featureless Al(Cr)2O3 scale formed over 

NiAl, albeit with occasional nodules of distinctive, rod-like needles, while an open, granular Cr2O3 scale 

formed over the Cr7C3. Unaffected regions away from the bond line exhibited similar granular Cr2O3

portions, presumably over the Cr7C3 impurity phases, with a very fine grain lenticular Al2O3 for the 

majority and remainder of the surface. On the LSHR side, the scale morphology was primarily nodular, 

having high Al, Cr levels. 
NASA/TM—2014-216679 17 

No evidence of interfacial damage or cracking was apparent, either at the exposed surfaces or 

polished cross-sections. A simple bilayer stress balance equation can be used to estimate the stress in the 

diffusion couple (Zhu & Miller, 1996): 

 

  s s c c c s

s c c s s c E t E t

TE E t

 Q   Q

D  D '

V 1 1

where the subscripts s and c refer to substrate (LSHR)and coating (Cr2AlC), respectively, and the other 

symbols, V, D, T, E, t, refer to stress, thermal expansion, temperature, Young’s modulus, and thickness, 

respectively. Using published values for these variables, the residual stress in the Cr2AlC can be 

calculated for cooling to 25 qC, assuming a stress free temperature at 800 qC, no edge effects and no 

bending. This stress was found to be 520 and 340 MPa, respectively, for the 0.3 and 1.5 mm Cr2AlC 

layers bonded to 1.8 mm of the LSHR alloy. This is compared to a flexural strength of 483 MPa measured 

for Cr2AlC (Wu-bian Tian, Wang, Zhang, Kan, & Li, 2007) and to multi-GPa levels predicted and 

measured for alumina scales on MCrAl alloys.(Lipkin & Clarke, 1996)(Tolpygo, Dryden, & Clarke, 

1998) This indicates that some yielding may be expected for thin Cr2AlC layers bonded to a superalloy, 

but would not approach the strain energy that induces spalling of alumina scales formed on Ni-alloys, 

albeit upon cooling from 1100 or 1200 qC. The oxidized edge exhibited no anomalous oxidation or 

synergistic degrading factor. The longer 1000 h test (17 cycles), using a 1.5 mm thick Cr2AlC wafer, 

appears to follow the same promising trend. 

Concluding Remarks 

Cr2AlC MAX phase is relatively stable with Ni-base superalloys in 800 qC cycling and may offer 

promise as a high temperature, strain tolerant, corrosion resistant coating. The mechanical stability 

exhibited here is at least an initial indication of a promising system. The use of 0.3 or 1.5 mm thick layers 

in a diffusion couple indicate a robustness not normally achieved for other oxidation-resistant coatings, 

which are generally limited to ~ 0.1 mm or less. Most of the interdiffusion resulted from hot pressing at 

1100 qC, with little growth observed at the expected 800 qC expected maximum service temperature. The 

ability of MAX phases to deform without cracking may provide some fatigue benefit as a coating. 

However the (brittle) NiAl diffusion zone may limit this benefit by its known fatigue debit as a coating 

for superalloys. The occurrence of some V-phase in the LSHR side is another cause of concern. It is not 

clear whether an interfacial region of fine alumina stringers is an artifact of the known Al2O3 impurity 

phase in the as-received Cr2AlC or a troublesome indigenous result of interdiffusion. In general, it is 

believed that lower processing temperatures will greatly decrease any detrimental diffusion effects. 

Oxidation of Cr2AlC is extremely slow at 800 qC and should pose no problem for surface 

degradation. As a carbide, oxygen diffusion through the MAX phase coating and oxygen-induced gas 

phase embrittlement (GPE) of the underlying alloy ought to be eliminated. As an alumina/chromia former 

with no Ni or Co exposed to the environment, Type II low temperature hot corrosion triggered by Ni-CoNa sulfate eutectics, is also expected to be minimal for Cr2AlC. 

Future Work 

It is recognized that numerous issues remain. Adequate Type II hot corrosion resistance must be fully 

demonstrated on bulk Cr2AlC. Quality coatings must be produced with good corrosion resistance and 

good diffusional phase stability, with no debit on fatigue life. In that regard, detrimental effects of the 

NiAl reaction zone, Cr7C3 depletion zone, and V formation must be evaluated in greater detail. 
NASA/TM—2014-216679 18 

References 

Barsoum, M. W., & El-RDJK\77KH0$;3KDVHVௗ8QLTXH1HZ&DUELGHDQG1LWULGH0DWHULDOV

American Scientist, 89 (July-August), 334–343. 

Birbilis, N., & Buchheit, R. G. (2008). Measurement and Discussion of Low-Temperature Hot Corrosion 

Damage Accumulation upon Nickel-Based Superalloy Rene 104. Metallurgical and Materials 

Transactions A, 39 (13), 3224–3232. doi:10.1007/s11661-008-9662-7 

Cruchley, S., Evans, H. E., Taylor, M. P., Hardy, M. C., & Stekovic, S. (2013). Chromia layer growth on 

a Ni-based superalloy: Sub-parabolic kinetics and the role of titanium. Corrosion Science, 75, 58–

66. doi:10.1016/j.corsci.2013.05.016 

Eklund, P., Beckers, M., Jansson, U., Högberg, H., & Hultman, L. (2010). The Mn+1AXn phases: 

Materials science and thin-film processing. Thin Solid Films, 518(8), 1851–1878. 

doi:10.1016/j.tsf.2009.07.184 

Encinas-Oropesa, A., Drew, G. L., Hardy, M. C., Leggett, A. J., Nicholls, J. R., & Simms, N. J. (2008). 

Effects of Oxidation and Hot Corrosion in a Nickel Disc Alloy. Superalloys 2008 (Eleventh 

International Symposium), 609–618. doi:10.7449/2008/Superalloys_2008_609_618 

Gabb, T. P., Gayda, J., Telesman, J., & Kantzos, P. T. (2005). Thermal and Mechanical Property 

Characterization of the Advanced Disk Alloy LSHR. NASA/TM—2005-213645, (June), 1–75. 

Gabb, T. P., & Miller, D. R. (2012). Formation of Minor Phases in a Nickel-Based Disk Superalloy. 

NASA/TM—2012-217604, (July), 1–31. 

Gabb, T. P., Sudbrack, C. K., Draper, S. L., MacKay, R. A., & Telesman, J. (2014). Effects of Long Term 

Exposures on Fatigue of PM Disk Superalloys. Materials Performance and Characterization, 3(2), 

1–24. doi:10.1520/MPC20130037 

Gulbiski, W. (2004). Ti-Si-C sputter deposited thin film coatings. Surface and Coatings Technology, 180-

181, 341–346. doi:10.1016/j.surfcoat.2003.10.084 

Hajas, D. E., to Baben, M., Hallstedt, B., Iskandar, R., Mayer, J., & Schneider, J. M. (2011). Oxidation of 

Cr2AlC coatings in the temperature range of 1230 to 1410 °C. Surface and Coatings Technology, 

206(4), 591–598. doi:10.1016/j.surfcoat.2011.03.086 

Karabela, A., Zhao, L. G., Lin, B., Tong, J., & Hardy, M. C. (2013). Oxygen diffusion and crack growth 

for a nickel-based superalloy under fatigue-oxidation conditions. Materials Science and 

Engineering: A, 567, 46–57. doi:10.1016/j.msea.2012.12.088 

Karabela, A., Zhao, L. G., Tong, J., Simms, N. J., Nicholls, J. R., & Hardy, M. C. (2011). Effects of 

cyclic stress and temperature on oxidation damage of a nickel-based superalloy. Materials Science 

and Engineering: A, 528(19-20), 6194–6202. doi:10.1016/j.msea.2011.04.029 

Kitaguchi, H. S., Li, H. Y., Evans, H. E., Ding, R. G., Jones, I. P., Baxter, G., & Bowen, P. (2013). 

Oxidation ahead of a crack tip in an advanced Ni-based superalloy. Acta Materialia, 61(6), 1968–

1981. doi:10.1016/j.actamat.2012.12.017 

Lee, D. B., Nguyen, T. D., Han, J. H., & Park, S. W. (2007). Oxidation of Cr2AlC at 1300°C in air. 

Corrosion Science, 49(10), 3926–3934. doi:10.1016/j.corsci.2007.03.044 

Lee, D. B., Nguyen, T. D., & Park, S. W. (2011). Corrosion of Cr2AlC in Ar/1%SO2 Gas Between 900 

and 1200 °C. Oxidation of Metals, 75(5-6), 313–323. doi:10.1007/s11085-011-9233-y 

Lee, D. B., & Park, S. W. (2011). Corrosion of Ti3AlC2 at 800–1100°C in Ar–0.2% SO2 gas atmosphere. 

Corrosion Science, 53(8), 2645–2650. doi:10.1016/j.corsci.2011.05.001 

Li, J. J., Li, M. S., Xiang, H. M., Lu, X. P., & Zhou, Y. C. (2011). Short-term oxidation resistance and 

degradation of Cr2AlC coating on M38G superalloy at 900–1100°C. Corrosion Science, 53(11), 

3813–3820. doi:10.1016/j.corsci.2011.07.032 

Lin, Z. J., Li, M. S., Wang, J. Y., & Zhou, Y. C. (2007). High-temperature oxidation and hot corrosion of 

Cr2AlC. Acta Materialia, 55(18), 6182–6191. doi:10.1016/j.actamat.2007.07.024 

Lin, Z., Zhou, Y., Li, M., & Wang, J. (2006a). Hot corrosion and protection of Ti2AlC against Na2SO4

salt in air. Journal of the European Ceramic Society, 26(16), 3871–3879. 

doi:10.1016/j.jeurceramsoc.2005.12.004 
NASA/TM—2014-216679 19 

Lin, Z., Zhou, Y., Li, M., & Wang, J. (2006b). Improving the Na2SO4-induced corrosion resistance of 

Ti3AlC2 by pre-oxidation in air. Corrosion Science, 48(10), 3271–3280. 

doi:10.1016/j.corsci.2005.11.005 

Lipkin, D. M., & Clarke, D. R. (1996). Measurement of the stress in oxide scales formed by oxidation of 

alumina-forming alloys. Oxidation of Metals, 45, 267–280. doi:10.1007/BF01046985 

Mercer, C., Kawagishi, K., Tomimatsu, T., Hovis, D., & Pollock, T. M. (2011). A comparative 

investigation of oxide formation on EQ (Equilibrium) and NiCoCrAlY bond coats under stepped 

thermal cycling. Surface and Coatings Technology, 205, 3066–3072. 

doi:10.1016/j.surfcoat.2010.11.026 

Moverare, J. J., & Johansson, S. (2010). Damage mechanisms of a high-Cr single crystal superalloy 

during thermomechanical fatigue. Materials Science and Engineering: A, 527(3), 553–558. 

doi:10.1016/j.msea.2009.08.023 

Sato, A., Chiu, Y.-L., & Reed, R. C. (2011). Oxidation of nickel-based single-crystal superalloys for 

industrial gas turbine applications. Acta Materialia, 59(1), 225–240. 

doi:10.1016/j.actamat.2010.09.027 

Sudbrack, C., Draper, S., Gorman, T., Telesman, J., Gabb, T., & Hull, D. (2012). Oxidation and the 

Effects of High Temperature Exposures on Notched Fatigue Life of an Advanced Powder 

Metallurgy Disk Superalloy. In J. T. E. Huron, R. Reed, M. Mills, R. Montero, P. Portella (Ed.), 

Superalloys 2012: 12th International Symposium on Superalloys (pp. 863–872). Seven Springs PA: 

TMS, Warrendale, PA. 

Sumner, J., Encinas-Oropesa, A., Simms, N. J., & Oakey, J. E. (2011). High temperature oxidation and 

corrosion of gas turbine component materials in burner rig exposures. Materials at High 

Temperatures, 28(4), 369–376. doi:10.3184/096034011X13198126967382 

Tallman, D. J., Anasori, B., & Barsoum, M. W. (2013). A Critical Review of the Oxidation of Ti2AlC, 

Ti3AlC2 and Cr2AlC in Air. Materials Research Letters, 1(3), 115–125. 

doi:10.1080/21663831.2013.806364 

Tian, W., Wang, P., Kan, Y., & Zhang, G. (2008). Oxidation behavior of Cr2AlC ceramics at 1,100 and 

1,250 °C. Journal of Materials Science, 43(8), 2785–2791. doi:10.1007/s10853-008-2516-2 

Tian, W., Wang, P., Zhang, G., Kan, Y., & Li, Y. (2007). Mechanical Properties of Cr2AlC Ceramics. 

Journal of the American Ceramic Society, 90(5), 1663–1666. doi:10.1111/j.1551-

2916.2007.01634.x 

Tolpygo, V. K., Dryden, J. R., 	&ODUNH'5'HWHUPLQDWLRQRIWKHJURZWKVWUHVVDQGVWUDLQLQĮAl2O3 scales during the oxidation of Fe–22Cr–4.8Al–0.3Y alloy. Acta Materialia, 46, 927–937. 

doi:10.1016/S1359-6454(97)00306-6 

Walter, C., Sigumonrong, D. P., El-Raghy, T., & Schneider, J. M. (2006). Towards large area deposition 

of Cr2AlC on steel. Thin Solid Films, 515(2), 389–393. doi:10.1016/j.tsf.2005.12.219 

Wang, Q. M., Flores Renteria, a., Schroeter, O., Mykhaylonka, R., Leyens, C., Garkas, W., & to Baben, 

M. (2010). Fabrication and oxidation behavior of Cr2AlC coating on Ti6242 alloy. Surface and 

Coatings Technology, 204(15), 2343–2352. doi:10.1016/j.surfcoat.2010.01.002 

Wang, Q. M., Mykhaylonka, R., Flores Renteria, a., Zhang, J. L., Leyens, C., & Kim, K. H. (2010). 

Improving the high-WHPSHUDWXUHR[LGDWLRQUHVLVWDQFHRIDȕ–Ȗ7L$ODOOR\E\D&U2AlC coating. 

Corrosion Science, 52(11), 3793–3802. doi:10.1016/j.corsci.2010.07.031 

Woodford, D. A. (2006). Gas phase embrittlement and time dependent cracking of nickel based 

superalloys. Energy Materials: Materials Science and Engineering for Energy Systems, 1, 59–79. 

doi:10.1179/174892306X99679 

Zhu, D., & Miller, R. A. (1996). Evaluation of Oxidation Damage in Thermal Barrier Systems. NASA 

TM–107360, ARL–TR–125, 1–20.