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Creep in directionally solidified NiAl–Mo eutectics
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Scripta Materialia 65 (2011) 699–702 www.elsevier.com/locate/scriptamat
Creep in directionally solidified NiAl–Mo eutectics M. Dudova´,a K. Kucharˇova´,a T. Barta´k,a H. Bei,b E.P. George,b,c Ch. Somsend and A. Dlouhy´a,⇑ a
Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Brno 616 62, Czech Republic b Materials Science & Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA c Materials Science & Engineering Department, University of Tennessee, Knoxville, TN 37996, USA d Institut fu¨r Werkstoffe, Ruhr-Universita¨t Bochum, D-44780 Bochum, Germany Received 25 April 2011; revised 25 June 2011; accepted 7 July 2011 Available online 14 July 2011
A directionally solidified NiAl–Mo eutectic and an NiAl intermetallic, having respective nominal compositions Ni–45.5Al–9Mo and 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 the plastically deforming matrix up to a critical stress of the order of 2.5 GPa, at which point they yield. Microstructural evidence is provided for the dislocation-mediated stress transfer from the NiAl to the Mo phase. Ó 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 composites exhibit microstructures of major (NiAl matrix) and minor (Mo fibre) eutectic phases. Recent advances in processing have resulted in NiAl–Mo microstructures of unprecedented quality in terms of alloy purity and microstructural regularity [1] (see Fig. 1). These new achievements call for new creep studies. It has long been known that composite creep behaviour is mostly controlled by interfaces separating the composite constituents. In particular, the matrix–reinforcement interactions can cause a qualitative difference between the creep characteristics of the composite and those of the matrix alloy. It has been shown that class I Al–Mg alloy can switch its behaviour to class II [2] on the introduction of strong alumina fibres [3]. Furthermore, the alumina reinforcement improves the creep strength of commercial Al-based alloys by decelerating the composite creep rate by several orders of magnitude [4]. It has been proposed that zones of a higher dislocation density (work-hardened zones; WHZs) form around the reinforcement phase [5,6]. The associated stress redistribution [7] hinders deformation in the matrix, facilitates recovery processes at the fibre ends and may cause damage to the reinforcement [5]. These scenarios were proposed for materials that, although technologically important, contained truly
⇑ Corresponding author. Tel.: +420 532 290 412; fax: +420 541 218 657; e-mail: [email protected]
complicated distributions of reinforcement phases [8,9]. Limited experimental data are available for simpler reinforcement–matrix geometries of in situ composites [10– 14], even though these systems are ideal for testing various modelling approaches. Based on Vickers hardness measurements, Ferrandini et al. [12] reported that the NiAl intermetallic and the NiAl–Mo eutectic possessed comparable strength at temperatures above 873 K. In contrast, high-temperature compression creep tests of NiAl–X eutectics (X = Cr, Mo) and the NiAl intermetallic showed that, under similar external conditions, the creep rate of the eutectics can be up to five orders of magnitude lower than that of the NiAl intermetallic [13,14]. The strengthening effect observed in these studies has not been fully clarified on a microstructural basis. It has been shown only recently that micro- to nanoscale reinforcement phases in in situ composite materials exhibit strengths that are in the range of the expected theoretical strength [15]. Since the reinforcement phase in composite materials takes on the load redistributed from the soft matrix, the new results would suggest that these reinforcement properties could account for the composite load-bearing capacity at high temperatures. Experimental alloys Ni–45.5Al–9Mo (NiAl–Mo in situ composite) and Ni–45.2Al (at.%) (NiAl matrix alloy) were melted in a high-vacuum arc melting furnace, starting with the elemental metals. After several flipping and remelting cycles, the ingots were drop cast into a
1359-6462/$ - see front matter Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2011.07.019
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Figure 1. SEM image showing regular distribution of fine Mo fibres in a (0 0 1) cross-section plane of the NiAl–Mo in situ composite. The fibres are oriented end-on.
copper mould. The ingots were subsequently remelted and directionally solidified in a high-temperature floating zone furnace at a growth rate of 60 mm h 1. The resulting NiAl–Mo cylindrical crystals exhibited a regular distribution (14% by volume) of long Mo alloy (Mo–10Al–4Ni, at.%) fibres inside the NiAl matrix. Fibre and matrix crystals were [0 0 1] – oriented along the growth direction. The fibres have roughly square or rectangular cross-sections, with typical edge lengths between 400 and 800 nm. Further details of the processing and initial microstructure can be found elsewhere [1]. Cylindrical compression creep specimens (height 12 mm and diameter 5 mm) were cut out of the directionally solidified crystals using a spark erosion cutter. The compression axis was parallel to the [0 0 1] growth direction. Constant applied stress creep tests were performed at temperatures in the range between 1073 and 1173 K in purified argon mixed with 5 vol.% hydrogen [16,17]. True strain–time readings were continuously recorded by a PC-based data acquisition system [18]. The testing temperature was maintained constant within ±1 K along the specimen gauge length. Metallographic cross-sections perpendicular to the [0 0 1] growth direction were studied by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Thin TEM foils were investigated using standard diffraction contrast techniques, including selected area diffraction. Furthermore, convergent beam diffraction and scanning TEM were combined with energy-dispersive spectroscopy (EDS) chemical analysis. Three microscopes, CM12, CM20 and Tecnai F20, operating at 120 and 200 kV, were employed. The as-solidified microstructure of the in situ NiAl–Mo composite is shown in Figure 1. The SEM image reveals a regular pattern of fine Mo fibres in the (0 0 1) cross-section plane. These microstructures were subjected to compression creep. Figure 2a illustrates typical deformation kinetics at 1073 K for the NiAl matrix alloy (blue curves, applied stresses 150 and 250 MPa) and for the NiAl–Mo in situ composite (red curve, applied stress 250 MPa). After a normal metal-type primary transition, which accounts for the first 4% of compression creep strain, the matrix NiAl alloy creeps in a steady-state regime for the next more than 20% strain up to when the tests were interrupted. In contrast, the creep strain accumulation kinetics of the NiAl–Mo composite are quite different. For similar external conditions, the composite creep rate
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 K and (b) 1173 K. (c) Minimum creep rates recorded at 1173 K vs. applied stress; literature data [20,21] for the creep rate of pure Mo are shown.
drops by several orders of magnitude and reaches a sharp minimum at about 1% strain. The creep rate then gradually increases with increasing strain and approaches a stationary value for strains higher than 10%. The difference between the composite minimum creep rate and the steady-state creep rate of the NiAl intermetallic at 1023 K and 250 MPa is more then five orders of magnitude. Similar deformation kinetics were observed at 1173 K, as shown in Figure 2b. Here, the difference in creep rates between the NiAl intermetallic and the NiAl–Mo eutectic exceeds six orders of magnitude. Figure 2c summarizes the minimum creep rates of the NiAl–Mo composite and the NiAl matrix at 1173 K for different applied stresses. The data obtained for the composite include results of two round robin tests performed at TU Darmstadt [19]. All composite data can be represented by a Norton law with a stress exponent n = 14. Figure 2c also presents literature data on the minimum creep rate of pure Mo measured at the slightly lower temperature of 1144 K [20,21]. Together,
M. Dudova´ et al. / Scripta Materialia 65 (2011) 699–702
the data shown in Figure 2c suggest that NiAl and Mo, when combined in the form of long Mo fibres embedded in an NiAl matrix, exhibit a composite creep strength that is orders of magnitude higher than that of either one alone. The NiAl–Mo microstructure was investigated by TEM before and after creep. The TEM bright-field image in Figure 3a illustrates the initial state of the composite before creep. In the transverse section, the oval-shaped dislocation-free Mo fibres are embedded in a dislocation-containing NiAl matrix. A matrix dislocation density is as high as 5.5 1013 m 2. We suggest that this dislocation density is due to matrix plasticity [22], which relaxes thermal mismatch strains during solidification of the NiAl–Mo composite [23,24]. Dislocations are distributed quite uniformly in the NiAl matrix, and neither accumulation of dislocations at the fibre matrix interfaces nor formation of subgrain boundaries is observed (see Fig. 3a). In contrast to the dislocation and particle-free fibres in the initial state, after 35% creep strain accumulated at 1173 K and 300 MPa, the fibres contain many dislocations and NiAl precipitates with characteristic sizes of 15 nm and smaller. A typical example is shown in Figure 3b, where a shrinking dislocation loop interacts with NiAl nanoparticles inside an Mo fibre. The nature of the NiAl nanoparticles was characterized by EDS using a fine TEM electron probe. Microstructural changes observed in the NiAl matrix after creep are documented in Figure 4a and b. During creep, dislocations in the NiAl matrix organize into low-angle boundaries. This is shown in Figure 4a after creep at 1173 K and 200 MPa, and even more clearly in Figure 4b after creep at the higher applied stress of 300 MPa. The low-angle dislocation boundaries interact strongly with the Mo fibres, and these interactions contribute to the transfer of stress to the fibres. These interactions may also initiate cutting events in which NiAl matrix dislocations enter Mo fibres. A detailed investigation of the cutting mechanism is in progress. It is also apparent, mainly in Figure 4a, that free dislocations inside subgrains accumulate at the fibre–matrix interfaces and form WHZs [4–6]. The zones extend up to a distance of 500 nm into the NiAl matrix. A preliminary assessment shows that, on approaching a fibre, the free dislocation density inside subgrains increases from an average value of 5.7 1013 m 2 to a local density of 1.77 1014 m 2 in the WHZs. The tendency for WHZ formation is less clear after creep at 300 MPa. In this case, the average free dislocation density between fibres reaches 1.14 1014 m 2 and a considerably denser system of subgrain boundaries is observed (see Fig. 4b). We note that similar WHZs were reported by Misra et al. after room temperature deformation [25]. Results obtained in the present study clearly show that the formidable high-temperature strength of NiAl–Mo in situ composite is associated with the presence of fine Mo fibres. Since fibres are essentially dislocation-free in the as-solidified microstructure, they can support the high stresses associated with dense dislocations structures of WHZs that form in the early stages of creep (see also [4,5]). The dislocation stress fields are fully taken on by Mo fibres up to the point when the fibres either yield [15] or are broken under the load
701
Figure 3. (a) Dislocation-free Mo fibres and thermally induced dislocations in NiAl matrix of NiAl–Mo composite before creep. (b) Dislocated Mo fibre in NiAl–Mo subjected to 35% compression creep strain at 1173 K and 300 MPa.
[4,5]. In the compression mode investigated here we expect that yielding of fibres due to dislocation cutting is more likely than other fibre damage mechanisms. Creep strain of the order of 1% is clearly sufficient to load the fibres to the critical point. This can be inferred from the creep strain accumulation kinetics recorded for the NiAl–Mo samples (see Fig. 2a and b), where creep strains of this magnitude characterize sharp creep rate minima. We suggest that these minima correspond to the load-bearing capacity of dislocation-free Mo fibres, that is, to a statistically significant yielding of fibres in the composite. A modelling study is required to rationalize the observed values of the stress exponent that describes the stress dependences of the minimum creep rate. On further straining beyond the creep rate minimum, the creeping zones in the fibre spread along the fibre length and the overall creep rate of the NiAl–Mo composite accelerates, eventually reaching a new stationary regime. It may be expected that this new stationary creep rate is controlled by deformation processes in the Mo phase, which, in the investigated range of external conditions, exhibits higher creep strength than the intermetallic NiAl phase (see Fig. 2c). At these advanced stages of compression creep, contributions to the creep strain accumulation kinetics due to the fibre breakage cannot be ruled out.
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(5) Using a room-temperature elastic modulus of 270 GPa for the Mo phase [1] and the strains attained at the creep rate minima, dislocationfree Mo fibres can support compressive stresses of the order of 2.5 GPa in the early stages of creep. (6) NiAl matrix dislocations arrange into low-angle boundaries and form work-hardened zones around Mo fibres during creep. These dislocation structures mediate the stress transfer from the matrix to the fibres [4,5]. Two NiAl–Mo creep tests at 1173 K were performed under the supervision of Prof. M. Heilmaier using the facilities of TU Darmstadt. The final electropolishing of TEM foils was performed by Dipl.-Ing. T. Simon at RU Bochum. Financial support was received from the Czech Science Foundation, Contract number 202/09/2073. H.B. and E.P.G. were supported by the US Department of Energy, Basic Energy Sciences, Materials Sciences and Engineering Division.
Figure 4. Interactions between Mo fibres, subgrain boundaries and free dislocations inside subgrains documented after compression creep (a) at 1173 K and 200 MPa, accumulated strain 33% and (b) at 1173 K and 300 MPa, accumulated strain 35%.
In summary, [0 0 1] -oriented Ni–45.2Al intermetallic crystals (matrix alloy) and [0 0 1]-oriented Ni–45.5Al–9Mo (all at.%) eutectics (in situ composites) were tested in a compression creep regime at temperatures between 1073 and 1173 K, and associated microstructural changes were investigated using SEM and TEM techniques. Based on the obtained experimental results, we can draw the following conclusions: (1) While Ni–45.2Al matrix crystals show metaltype creep behaviour, the Ni–45.5Al–9Mo in situ composites exhibit a sharp creep rate minimum at strains of 1–2%. (2) At the investigated temperatures and applied stresses, the Ni–45.2Al matrix crystals and the bulk Mo samples [20,21] creep considerably faster than the Ni–45.5Al–9Mo in situ composite. (3) Fine Mo fibres (typical diameter 600 nm) regularly distributed in the Ni–45.5Al–9Mo in situ composite result in minimum creep rates that are up to seven orders of magnitude lower than the corresponding minimum creep rates of the Ni–45.2Al matrix alloy. (4) Stress dependence of the composite minimum creep rate can be phenomenologically represented by a Norton law with stress exponent n = 14.
[1] H. Bei, E.P. George, Acta Mater. 53 (2005) 69. ˇ adek, Creep in Metallic Materials, Elsevier, Amsterdam, [2] J. C 1988. [3] G. Kaustrater, B. Skrotzki, G. Eggeler, Mater. Sci. Eng. A 319 (2001) 716. [4] A. Dlouhy´, N. Merk, G. Eggeler, Acta Metall. Mater. 41 (1993) 3245. [5] A. Dlouhy´, G. Eggeler, N. Merk, Acta Metall. Mater. 43 (1995) 535. [6] G. Eggeler, J.C. Earthman, Scr. Mater. 38 (1997) 341. [7] H. Mughrabi, Mater. Sci. Eng. A 85 (1987) 15. [8] A. Dlouhy´, G. Eggeler, Pract. Metallography 30 (1993) 172. [9] G. Garces, G. Bruno, A. Wanner, Acta Mater. 55 (2007) 5389. [10] J. Lapin, L. Ryelandt, F. Delannay, J. Phys. IV 3 (1993) 1753. [11] J. Lapin, F. Delannay, Metall. Mater. Trans. A 26 (1995) 2053. [12] P. Ferrandini, W.W. Batista, R. Caram, J. Alloys Comp. 381 (2004) 91. [13] D.R. Johnson, X.F. Chen, B.F. Oliver, R.D. Noebe, J.D. Whittenberger, Intermetallics 3 (1995) 99. [14] T. Haenschke, A. Gali, M. Heilmaier, M. Kru¨ger, H. Bei, E.P. George, J. Phys.: Conf. Ser. 240 (2010) 012063. [15] H. Bei, S. Shim, E.P. George, M.K. Miller, E.G. Herbert, G.M. Pharr, Scr. Mater. 57 (2007) 397. ˇ adek, J. Test. Eval. 4 (1976) 26. [16] T. Hostinsky´, J. C ˇ adek, Met. Mater. 24 (1986) 599. [17] F. Dobesˇ, O. Zveˇrˇina, J. C [18] K. Milicˇka, Internal Report, IPM, Brno, 1993. [19] M. Heilmaier, personal communication, 2010. [20] J.W. Pugh, Trans. ASM 47 (1955) 984. [21] H.J. Frost, M.F. Ashby, Deformation-Mechanism Maps, Pergamon Press, Oxford, 1982. [22] M.J. Mills, R. Srinivasan, M.S. Daw, Phil. Mag. A 77 (1998) 801. [23] H. Bei, E.P. George, D.W. Brown, G.M. Pharr, H. Choo, W.D. Porter, M.A.M. Bourke, J. Appl. Phys. 97 (2005) 123503–123511. [24] R.I. Barabash, H. Bei, Y. Gao, G.E. Ice, E.P. George, J. Mater. Res. 25 (2010) 199. [25] A. Misra, Z.L. Wu, M.T. Kush, R. Gibala, Mater. Sci. Eng. A 239–240 (1997) 75.
Scripta Materialia 65 (2011) 699–702 www.elsevier.com/locate/scriptamat
Creep in directionally solidified NiAl–Mo eutectics M. Dudova´,a K. Kucharˇova´,a T. Barta´k,a H. Bei,b E.P. George,b,c Ch. Somsend and A. Dlouhy´a,⇑ a
Institute of Physics of Materials, Academy of Sciences of the Czech Republic, Brno 616 62, Czech Republic b Materials Science & Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA c Materials Science & Engineering Department, University of Tennessee, Knoxville, TN 37996, USA d Institut fu¨r Werkstoffe, Ruhr-Universita¨t Bochum, D-44780 Bochum, Germany Received 25 April 2011; revised 25 June 2011; accepted 7 July 2011 Available online 14 July 2011
A directionally solidified NiAl–Mo eutectic and an NiAl intermetallic, having respective nominal compositions Ni–45.5Al–9Mo and 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 the plastically deforming matrix up to a critical stress of the order of 2.5 GPa, at which point they yield. Microstructural evidence is provided for the dislocation-mediated stress transfer from the NiAl to the Mo phase. Ó 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 composites exhibit microstructures of major (NiAl matrix) and minor (Mo fibre) eutectic phases. Recent advances in processing have resulted in NiAl–Mo microstructures of unprecedented quality in terms of alloy purity and microstructural regularity [1] (see Fig. 1). These new achievements call for new creep studies. It has long been known that composite creep behaviour is mostly controlled by interfaces separating the composite constituents. In particular, the matrix–reinforcement interactions can cause a qualitative difference between the creep characteristics of the composite and those of the matrix alloy. It has been shown that class I Al–Mg alloy can switch its behaviour to class II [2] on the introduction of strong alumina fibres [3]. Furthermore, the alumina reinforcement improves the creep strength of commercial Al-based alloys by decelerating the composite creep rate by several orders of magnitude [4]. It has been proposed that zones of a higher dislocation density (work-hardened zones; WHZs) form around the reinforcement phase [5,6]. The associated stress redistribution [7] hinders deformation in the matrix, facilitates recovery processes at the fibre ends and may cause damage to the reinforcement [5]. These scenarios were proposed for materials that, although technologically important, contained truly
⇑ Corresponding author. Tel.: +420 532 290 412; fax: +420 541 218 657; e-mail: [email protected]
complicated distributions of reinforcement phases [8,9]. Limited experimental data are available for simpler reinforcement–matrix geometries of in situ composites [10– 14], even though these systems are ideal for testing various modelling approaches. Based on Vickers hardness measurements, Ferrandini et al. [12] reported that the NiAl intermetallic and the NiAl–Mo eutectic possessed comparable strength at temperatures above 873 K. In contrast, high-temperature compression creep tests of NiAl–X eutectics (X = Cr, Mo) and the NiAl intermetallic showed that, under similar external conditions, the creep rate of the eutectics can be up to five orders of magnitude lower than that of the NiAl intermetallic [13,14]. The strengthening effect observed in these studies has not been fully clarified on a microstructural basis. It has been shown only recently that micro- to nanoscale reinforcement phases in in situ composite materials exhibit strengths that are in the range of the expected theoretical strength [15]. Since the reinforcement phase in composite materials takes on the load redistributed from the soft matrix, the new results would suggest that these reinforcement properties could account for the composite load-bearing capacity at high temperatures. Experimental alloys Ni–45.5Al–9Mo (NiAl–Mo in situ composite) and Ni–45.2Al (at.%) (NiAl matrix alloy) were melted in a high-vacuum arc melting furnace, starting with the elemental metals. After several flipping and remelting cycles, the ingots were drop cast into a
1359-6462/$ - see front matter Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2011.07.019
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M. Dudova´ et al. / Scripta Materialia 65 (2011) 699–702
Figure 1. SEM image showing regular distribution of fine Mo fibres in a (0 0 1) cross-section plane of the NiAl–Mo in situ composite. The fibres are oriented end-on.
copper mould. The ingots were subsequently remelted and directionally solidified in a high-temperature floating zone furnace at a growth rate of 60 mm h 1. The resulting NiAl–Mo cylindrical crystals exhibited a regular distribution (14% by volume) of long Mo alloy (Mo–10Al–4Ni, at.%) fibres inside the NiAl matrix. Fibre and matrix crystals were [0 0 1] – oriented along the growth direction. The fibres have roughly square or rectangular cross-sections, with typical edge lengths between 400 and 800 nm. Further details of the processing and initial microstructure can be found elsewhere [1]. Cylindrical compression creep specimens (height 12 mm and diameter 5 mm) were cut out of the directionally solidified crystals using a spark erosion cutter. The compression axis was parallel to the [0 0 1] growth direction. Constant applied stress creep tests were performed at temperatures in the range between 1073 and 1173 K in purified argon mixed with 5 vol.% hydrogen [16,17]. True strain–time readings were continuously recorded by a PC-based data acquisition system [18]. The testing temperature was maintained constant within ±1 K along the specimen gauge length. Metallographic cross-sections perpendicular to the [0 0 1] growth direction were studied by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Thin TEM foils were investigated using standard diffraction contrast techniques, including selected area diffraction. Furthermore, convergent beam diffraction and scanning TEM were combined with energy-dispersive spectroscopy (EDS) chemical analysis. Three microscopes, CM12, CM20 and Tecnai F20, operating at 120 and 200 kV, were employed. The as-solidified microstructure of the in situ NiAl–Mo composite is shown in Figure 1. The SEM image reveals a regular pattern of fine Mo fibres in the (0 0 1) cross-section plane. These microstructures were subjected to compression creep. Figure 2a illustrates typical deformation kinetics at 1073 K for the NiAl matrix alloy (blue curves, applied stresses 150 and 250 MPa) and for the NiAl–Mo in situ composite (red curve, applied stress 250 MPa). After a normal metal-type primary transition, which accounts for the first 4% of compression creep strain, the matrix NiAl alloy creeps in a steady-state regime for the next more than 20% strain up to when the tests were interrupted. In contrast, the creep strain accumulation kinetics of the NiAl–Mo composite are quite different. For similar external conditions, the composite creep rate
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 K and (b) 1173 K. (c) Minimum creep rates recorded at 1173 K vs. applied stress; literature data [20,21] for the creep rate of pure Mo are shown.
drops by several orders of magnitude and reaches a sharp minimum at about 1% strain. The creep rate then gradually increases with increasing strain and approaches a stationary value for strains higher than 10%. The difference between the composite minimum creep rate and the steady-state creep rate of the NiAl intermetallic at 1023 K and 250 MPa is more then five orders of magnitude. Similar deformation kinetics were observed at 1173 K, as shown in Figure 2b. Here, the difference in creep rates between the NiAl intermetallic and the NiAl–Mo eutectic exceeds six orders of magnitude. Figure 2c summarizes the minimum creep rates of the NiAl–Mo composite and the NiAl matrix at 1173 K for different applied stresses. The data obtained for the composite include results of two round robin tests performed at TU Darmstadt [19]. All composite data can be represented by a Norton law with a stress exponent n = 14. Figure 2c also presents literature data on the minimum creep rate of pure Mo measured at the slightly lower temperature of 1144 K [20,21]. Together,
M. Dudova´ et al. / Scripta Materialia 65 (2011) 699–702
the data shown in Figure 2c suggest that NiAl and Mo, when combined in the form of long Mo fibres embedded in an NiAl matrix, exhibit a composite creep strength that is orders of magnitude higher than that of either one alone. The NiAl–Mo microstructure was investigated by TEM before and after creep. The TEM bright-field image in Figure 3a illustrates the initial state of the composite before creep. In the transverse section, the oval-shaped dislocation-free Mo fibres are embedded in a dislocation-containing NiAl matrix. A matrix dislocation density is as high as 5.5 1013 m 2. We suggest that this dislocation density is due to matrix plasticity [22], which relaxes thermal mismatch strains during solidification of the NiAl–Mo composite [23,24]. Dislocations are distributed quite uniformly in the NiAl matrix, and neither accumulation of dislocations at the fibre matrix interfaces nor formation of subgrain boundaries is observed (see Fig. 3a). In contrast to the dislocation and particle-free fibres in the initial state, after 35% creep strain accumulated at 1173 K and 300 MPa, the fibres contain many dislocations and NiAl precipitates with characteristic sizes of 15 nm and smaller. A typical example is shown in Figure 3b, where a shrinking dislocation loop interacts with NiAl nanoparticles inside an Mo fibre. The nature of the NiAl nanoparticles was characterized by EDS using a fine TEM electron probe. Microstructural changes observed in the NiAl matrix after creep are documented in Figure 4a and b. During creep, dislocations in the NiAl matrix organize into low-angle boundaries. This is shown in Figure 4a after creep at 1173 K and 200 MPa, and even more clearly in Figure 4b after creep at the higher applied stress of 300 MPa. The low-angle dislocation boundaries interact strongly with the Mo fibres, and these interactions contribute to the transfer of stress to the fibres. These interactions may also initiate cutting events in which NiAl matrix dislocations enter Mo fibres. A detailed investigation of the cutting mechanism is in progress. It is also apparent, mainly in Figure 4a, that free dislocations inside subgrains accumulate at the fibre–matrix interfaces and form WHZs [4–6]. The zones extend up to a distance of 500 nm into the NiAl matrix. A preliminary assessment shows that, on approaching a fibre, the free dislocation density inside subgrains increases from an average value of 5.7 1013 m 2 to a local density of 1.77 1014 m 2 in the WHZs. The tendency for WHZ formation is less clear after creep at 300 MPa. In this case, the average free dislocation density between fibres reaches 1.14 1014 m 2 and a considerably denser system of subgrain boundaries is observed (see Fig. 4b). We note that similar WHZs were reported by Misra et al. after room temperature deformation [25]. Results obtained in the present study clearly show that the formidable high-temperature strength of NiAl–Mo in situ composite is associated with the presence of fine Mo fibres. Since fibres are essentially dislocation-free in the as-solidified microstructure, they can support the high stresses associated with dense dislocations structures of WHZs that form in the early stages of creep (see also [4,5]). The dislocation stress fields are fully taken on by Mo fibres up to the point when the fibres either yield [15] or are broken under the load
701
Figure 3. (a) Dislocation-free Mo fibres and thermally induced dislocations in NiAl matrix of NiAl–Mo composite before creep. (b) Dislocated Mo fibre in NiAl–Mo subjected to 35% compression creep strain at 1173 K and 300 MPa.
[4,5]. In the compression mode investigated here we expect that yielding of fibres due to dislocation cutting is more likely than other fibre damage mechanisms. Creep strain of the order of 1% is clearly sufficient to load the fibres to the critical point. This can be inferred from the creep strain accumulation kinetics recorded for the NiAl–Mo samples (see Fig. 2a and b), where creep strains of this magnitude characterize sharp creep rate minima. We suggest that these minima correspond to the load-bearing capacity of dislocation-free Mo fibres, that is, to a statistically significant yielding of fibres in the composite. A modelling study is required to rationalize the observed values of the stress exponent that describes the stress dependences of the minimum creep rate. On further straining beyond the creep rate minimum, the creeping zones in the fibre spread along the fibre length and the overall creep rate of the NiAl–Mo composite accelerates, eventually reaching a new stationary regime. It may be expected that this new stationary creep rate is controlled by deformation processes in the Mo phase, which, in the investigated range of external conditions, exhibits higher creep strength than the intermetallic NiAl phase (see Fig. 2c). At these advanced stages of compression creep, contributions to the creep strain accumulation kinetics due to the fibre breakage cannot be ruled out.
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M. Dudova´ et al. / Scripta Materialia 65 (2011) 699–702
(5) Using a room-temperature elastic modulus of 270 GPa for the Mo phase [1] and the strains attained at the creep rate minima, dislocationfree Mo fibres can support compressive stresses of the order of 2.5 GPa in the early stages of creep. (6) NiAl matrix dislocations arrange into low-angle boundaries and form work-hardened zones around Mo fibres during creep. These dislocation structures mediate the stress transfer from the matrix to the fibres [4,5]. Two NiAl–Mo creep tests at 1173 K were performed under the supervision of Prof. M. Heilmaier using the facilities of TU Darmstadt. The final electropolishing of TEM foils was performed by Dipl.-Ing. T. Simon at RU Bochum. Financial support was received from the Czech Science Foundation, Contract number 202/09/2073. H.B. and E.P.G. were supported by the US Department of Energy, Basic Energy Sciences, Materials Sciences and Engineering Division.
Figure 4. Interactions between Mo fibres, subgrain boundaries and free dislocations inside subgrains documented after compression creep (a) at 1173 K and 200 MPa, accumulated strain 33% and (b) at 1173 K and 300 MPa, accumulated strain 35%.
In summary, [0 0 1] -oriented Ni–45.2Al intermetallic crystals (matrix alloy) and [0 0 1]-oriented Ni–45.5Al–9Mo (all at.%) eutectics (in situ composites) were tested in a compression creep regime at temperatures between 1073 and 1173 K, and associated microstructural changes were investigated using SEM and TEM techniques. Based on the obtained experimental results, we can draw the following conclusions: (1) While Ni–45.2Al matrix crystals show metaltype creep behaviour, the Ni–45.5Al–9Mo in situ composites exhibit a sharp creep rate minimum at strains of 1–2%. (2) At the investigated temperatures and applied stresses, the Ni–45.2Al matrix crystals and the bulk Mo samples [20,21] creep considerably faster than the Ni–45.5Al–9Mo in situ composite. (3) Fine Mo fibres (typical diameter 600 nm) regularly distributed in the Ni–45.5Al–9Mo in situ composite result in minimum creep rates that are up to seven orders of magnitude lower than the corresponding minimum creep rates of the Ni–45.2Al matrix alloy. (4) Stress dependence of the composite minimum creep rate can be phenomenologically represented by a Norton law with stress exponent n = 14.
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