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Scripta Materialia 65 (2011) 699–702

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Creep in directionally solidified NiAl–Mo eutectics

M. Dudova,a K. Kucharova,a T. Bartak,a H. Bei,b E.P. George,b,c Ch. Somsend andA. Dlouhya,⇑

aInstitute of Physics of Materials, Academy of Sciences of the Czech Republic, Brno 616 62, Czech RepublicbMaterials Science & Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA

cMaterials Science & Engineering Department, University of Tennessee, Knoxville, TN 37996, USAdInstitut fur Werkstoffe, Ruhr-Universitat Bochum, D-44780 Bochum, Germany

Received 25 April 2011; revised 25 June 2011; accepted 7 July 2011Available online 14 July 2011

A directionally solidified NiAl–Mo eutectic and an NiAl intermetallic, having respective nominal compositions Ni–45.5Al–9Moand Ni–45.2Al (at.%), were loaded in compression at 1073 and 1173 K. Formidable strengthening by regularly distributed Mo fibres(average diameter 600 nm, volume fraction 14%) was observed. The fibres can support compression stresses transferred from theplastically deforming matrix up to a critical stress of the order of 2.5 GPa, at which point they yield. Microstructural evidence isprovided for the dislocation-mediated stress transfer from the NiAl to the Mo phase.� 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: High-temperature creep; NiAl–Mo eutectic; Mo fibres; Directional solidification; Transmission electron microscopy

Directionally solidified NiAl–Mo compositesexhibit microstructures of major (NiAl matrix) and minor(Mo fibre) eutectic phases. Recent advances in process-ing have resulted in NiAl–Mo microstructures ofunprecedented quality in terms of alloy purity and micro-structural regularity [1] (see Fig. 1). These new achieve-ments call for new creep studies. It has long beenknown that composite creep behaviour is mostlycontrolled by interfaces separating the composite constit-uents. In particular, the matrix–reinforcement interac-tions can cause a qualitative difference between thecreep characteristics of the composite and those of thematrix alloy. It has been shown that class I Al–Mg alloycan switch its behaviour to class II [2] on the introductionof strong alumina fibres [3]. Furthermore, the aluminareinforcement improves the creep strength of commercialAl-based alloys by decelerating the composite creep rateby several orders of magnitude [4]. It has been proposedthat zones of a higher dislocation density (work-hardenedzones; WHZs) form around the reinforcement phase [5,6].The associated stress redistribution [7] hinders deforma-tion in the matrix, facilitates recovery processes at the fi-bre ends and may cause damage to the reinforcement [5].These scenarios were proposed for materials that,although technologically important, contained truly

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complicated distributions of reinforcement phases [8,9].Limited experimental data are available for simpler rein-forcement–matrix geometries of in situ composites [10–14], even though these systems are ideal for testing variousmodelling approaches. Based on Vickers hardness mea-surements, Ferrandini et al. [12] reported that the NiAlintermetallic and the NiAl–Mo eutectic possessed compa-rable strength at temperatures above 873 K. In contrast,high-temperature compression creep tests of NiAl–Xeutectics (X = Cr, Mo) and the NiAl intermetallic showedthat, under similar external conditions, the creep rate ofthe eutectics can be up to five orders of magnitude lowerthan that of the NiAl intermetallic [13,14]. The strength-ening effect observed in these studies has not been fullyclarified on a microstructural basis. It has been shownonly recently that micro- to nanoscale reinforcementphases in in situ composite materials exhibit strengthsthat are in the range of the expected theoretical strength[15]. Since the reinforcement phase in composite materialstakes on the load redistributed from the soft matrix, thenew results would suggest that these reinforcement prop-erties could account for the composite load-bearingcapacity at high temperatures.

Experimental alloys Ni–45.5Al–9Mo (NiAl–Moin situ composite) and Ni–45.2Al (at.%) (NiAl matrix al-loy) were melted in a high-vacuum arc melting furnace,starting with the elemental metals. After several flippingand remelting cycles, the ingots were drop cast into a

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Figure 2. Compression creep curves of the [0 0 1]-oriented NiAl and[0 0 1]-oriented NiAl–Mo for different applied stresses at (a) 1073 Kand (b) 1173 K. (c) Minimum creep rates recorded at 1173 K vs. appliedstress; literature data [20,21] for the creep rate of pure Mo are shown.

Figure 1. SEM image showing regular distribution of fine Mo fibres ina (0 0 1) cross-section plane of the NiAl–Mo in situ composite. Thefibres are oriented end-on.

700 M. Dudova et al. / Scripta Materialia 65 (2011) 699–702

copper mould. The ingots were subsequently remeltedand directionally solidified in a high-temperature float-ing zone furnace at a growth rate of 60 mm h�1. Theresulting NiAl–Mo cylindrical crystals exhibited a regu-lar distribution (14% by volume) of long Mo alloy(Mo–10Al–4Ni, at.%) fibres inside the NiAl matrix. Fi-bre and matrix crystals were [0 0 1] – oriented alongthe growth direction. The fibres have roughly squareor rectangular cross-sections, with typical edge lengthsbetween 400 and 800 nm. Further details of the process-ing and initial microstructure can be found elsewhere [1].

Cylindrical compression creep specimens (height12 mm and diameter 5 mm) were cut out of the direc-tionally solidified crystals using a spark erosion cutter.The compression axis was parallel to the [0 0 1] growthdirection. Constant applied stress creep tests were per-formed at temperatures in the range between 1073 and1173 K in purified argon mixed with 5 vol.% hydrogen[16,17]. True strain–time readings were continuously re-corded by a PC-based data acquisition system [18]. Thetesting temperature was maintained constant within±1 K along the specimen gauge length. Metallographiccross-sections perpendicular to the [0 0 1] growth direc-tion were studied by scanning electron microscopy(SEM) and transmission electron microscopy (TEM).Thin TEM foils were investigated using standard diffrac-tion contrast techniques, including selected area diffrac-tion. Furthermore, convergent beam diffraction andscanning TEM were combined with energy-dispersivespectroscopy (EDS) chemical analysis. Three micro-scopes, CM12, CM20 and Tecnai F20, operating at120 and 200 kV, were employed.

The as-solidified microstructure of the in situNiAl–Mo composite is shown in Figure 1. The SEM im-age reveals a regular pattern of fine Mo fibres in the (0 0 1)cross-section plane. These microstructures were subjectedto compression creep.

Figure 2a illustrates typical deformation kinetics at1073 K for the NiAl matrix alloy (blue curves, appliedstresses 150 and 250 MPa) and for the NiAl–Moin situ composite (red curve, applied stress 250 MPa).After a normal metal-type primary transition, which ac-counts for the first 4% of compression creep strain, thematrix NiAl alloy creeps in a steady-state regime forthe next more than 20% strain up to when the tests wereinterrupted. In contrast, the creep strain accumulationkinetics of the NiAl–Mo composite are quite different.For similar external conditions, the composite creep rate

drops by several orders of magnitude and reaches asharp minimum at about 1% strain. The creep rate thengradually increases with increasing strain and ap-proaches a stationary value for strains higher than10%. The difference between the composite minimumcreep rate and the steady-state creep rate of the NiAlintermetallic at 1023 K and 250 MPa is more then fiveorders of magnitude. Similar deformation kinetics wereobserved at 1173 K, as shown in Figure 2b. Here, thedifference in creep rates between the NiAl intermetallicand the NiAl–Mo eutectic exceeds six orders of magni-tude. Figure 2c summarizes the minimum creep ratesof the NiAl–Mo composite and the NiAl matrix at1173 K for different applied stresses. The data obtainedfor the composite include results of two round robintests performed at TU Darmstadt [19]. All compositedata can be represented by a Norton law with a stressexponent n = 14. Figure 2c also presents literature dataon the minimum creep rate of pure Mo measured at theslightly lower temperature of 1144 K [20,21]. Together,

Figure 3. (a) Dislocation-free Mo fibres and thermally induceddislocations in NiAl matrix of NiAl–Mo composite before creep. (b)Dislocated Mo fibre in NiAl–Mo subjected to 35% compression creepstrain at 1173 K and 300 MPa.

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the data shown in Figure 2c suggest that NiAl and Mo,when combined in the form of long Mo fibres embeddedin an NiAl matrix, exhibit a composite creep strengththat is orders of magnitude higher than that of eitherone alone.

The NiAl–Mo microstructure was investigated byTEM before and after creep. The TEM bright-field imagein Figure 3a illustrates the initial state of the compositebefore creep. In the transverse section, the oval-shapeddislocation-free Mo fibres are embedded in a disloca-tion-containing NiAl matrix. A matrix dislocation den-sity is as high as 5.5 � 1013 m�2. We suggest that thisdislocation density is due to matrix plasticity [22], whichrelaxes thermal mismatch strains during solidification ofthe NiAl–Mo composite [23,24]. Dislocations are distrib-uted quite uniformly in the NiAl matrix, and neither accu-mulation of dislocations at the fibre matrix interfaces norformation of subgrain boundaries is observed (seeFig. 3a). In contrast to the dislocation and particle-free fi-bres in the initial state, after 35% creep strain accumu-lated at 1173 K and 300 MPa, the fibres contain manydislocations and NiAl precipitates with characteristicsizes of 15 nm and smaller. A typical example is shownin Figure 3b, where a shrinking dislocation loop interactswith NiAl nanoparticles inside an Mo fibre.

The nature of the NiAl nanoparticles was character-ized by EDS using a fine TEM electron probe. Micro-structural changes observed in the NiAl matrix aftercreep are documented in Figure 4a and b. During creep,dislocations in the NiAl matrix organize into low-angleboundaries. This is shown in Figure 4a after creep at1173 K and 200 MPa, and even more clearly in Figure4b after creep at the higher applied stress of 300 MPa.The low-angle dislocation boundaries interact stronglywith the Mo fibres, and these interactions contributeto the transfer of stress to the fibres. These interactionsmay also initiate cutting events in which NiAl matrixdislocations enter Mo fibres. A detailed investigationof the cutting mechanism is in progress. It is also appar-ent, mainly in Figure 4a, that free dislocations insidesubgrains accumulate at the fibre–matrix interfacesand form WHZs [4–6]. The zones extend up to a dis-tance of 500 nm into the NiAl matrix. A preliminaryassessment shows that, on approaching a fibre, the freedislocation density inside subgrains increases from anaverage value of 5.7 � 1013 m�2 to a local density of1.77 � 1014 m�2 in the WHZs. The tendency for WHZformation is less clear after creep at 300 MPa. In thiscase, the average free dislocation density between fibresreaches 1.14 � 1014 m�2 and a considerably denser sys-tem of subgrain boundaries is observed (see Fig. 4b).We note that similar WHZs were reported by Misraet al. after room temperature deformation [25].

Results obtained in the present study clearly showthat the formidable high-temperature strength ofNiAl–Mo in situ composite is associated with the pres-ence of fine Mo fibres. Since fibres are essentially dislo-cation-free in the as-solidified microstructure, they cansupport the high stresses associated with dense disloca-tions structures of WHZs that form in the early stagesof creep (see also [4,5]). The dislocation stress fieldsare fully taken on by Mo fibres up to the point whenthe fibres either yield [15] or are broken under the load

[4,5]. In the compression mode investigated here we ex-pect that yielding of fibres due to dislocation cutting ismore likely than other fibre damage mechanisms. Creepstrain of the order of 1% is clearly sufficient to load thefibres to the critical point. This can be inferred from thecreep strain accumulation kinetics recorded for theNiAl–Mo samples (see Fig. 2a and b), where creepstrains of this magnitude characterize sharp creep rateminima. We suggest that these minima correspond tothe load-bearing capacity of dislocation-free Mo fibres,that is, to a statistically significant yielding of fibres inthe composite. A modelling study is required to rational-ize the observed values of the stress exponent that de-scribes the stress dependences of the minimum creeprate. On further straining beyond the creep rate mini-mum, the creeping zones in the fibre spread along the fi-bre length and the overall creep rate of the NiAl–Mocomposite accelerates, eventually reaching a new sta-tionary regime. It may be expected that this new station-ary creep rate is controlled by deformation processes inthe Mo phase, which, in the investigated range of exter-nal conditions, exhibits higher creep strength than theintermetallic NiAl phase (see Fig. 2c). At these advancedstages of compression creep, contributions to the creepstrain accumulation kinetics due to the fibre breakagecannot be ruled out.

Figure 4. Interactions between Mo fibres, subgrain boundaries andfree dislocations inside subgrains documented after compression creep(a) at 1173 K and 200 MPa, accumulated strain 33% and (b) at 1173 Kand 300 MPa, accumulated strain 35%.

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In summary, [0 0 1] -oriented Ni–45.2Al intermetal-lic crystals (matrix alloy) and [0 0 1]-orientedNi–45.5Al–9Mo (all at.%) eutectics (in situ composites)were tested in a compression creep regime at tempera-tures between 1073 and 1173 K, and associatedmicrostructural changes were investigated using SEMand TEM techniques. Based on the obtained experi-mental results, we can draw the following conclusions:

(1) While Ni–45.2Al matrix crystals show metal-type creep behaviour, the Ni–45.5Al–9Moin situ composites exhibit a sharp creep rateminimum at strains of 1–2%.

(2) At the investigated temperatures and appliedstresses, the Ni–45.2Al matrix crystals and thebulk Mo samples [20,21] creep considerably fas-ter than the Ni–45.5Al–9Mo in situ composite.

(3) Fine Mo fibres (typical diameter 600 nm) regu-larly distributed in the Ni–45.5Al–9Mo in situcomposite result in minimum creep rates thatare up to seven orders of magnitude lower thanthe corresponding minimum creep rates of theNi–45.2Al matrix alloy.

(4) Stress dependence of the composite minimumcreep rate can be phenomenologically repre-sented by a Norton law with stress exponentn = 14.

(5) Using a room-temperature elastic modulus of270 GPa for the Mo phase [1] and the strainsattained at the creep rate minima, dislocation-free Mo fibres can support compressive stressesof the order of 2.5 GPa in the early stages ofcreep.

(6) NiAl matrix dislocations arrange into low-angleboundaries and form work-hardened zonesaround Mo fibres during creep. These disloca-tion structures mediate the stress transfer fromthe matrix to the fibres [4,5].

Two NiAl–Mo creep tests at 1173 K were per-formed under the supervision of Prof. M. Heilmaierusing the facilities of TU Darmstadt. The final electro-polishing of TEM foils was performed by Dipl.-Ing. T.Simon at RU Bochum. Financial support was receivedfrom the Czech Science Foundation, Contract number202/09/2073. H.B. and E.P.G. were supported by theUS Department of Energy, Basic Energy Sciences,Materials Sciences and Engineering Division.

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