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Microstructural inhomogeneities introduced in a Zr-based bulk metallic glass upon low-temperature annealing
Materials Science and Engineering A 491 (2008) 124–130
Microstructural inhomogeneities introduced in a Zr-based bulk metallic glass upon low-temperature annealing N. Van Steenberge a,∗ , A. Concustell a , J. Sort b , J. Das c , N. Mattern c , A. Gebert c , S. Suri˜nach a , J. Eckert c , M.D. Bar´o a b
a Departament de F´ısica, Universitat Aut` onoma de Barcelona, 08913 Bellaterra, Spain Instituci´o Catalana de Recerca i Estudis Avan¸cats and Departament de F´ısica, Universitat Aut`onoma de Barcelona, 08193 Bellaterra, Spain c Leibniz-Institut f¨ ur Festk¨orper- und Werkstoffforschung Dresden, P.O. Box 27 00 16, 01171 Dresden, Germany
Received 19 October 2007; received in revised form 21 January 2008; accepted 25 January 2008
Abstract Due to their exceptionally high yield strength and yield strain as compared to conventional metallic materials, bulk metallic glasses (BMGs) represent a class of promising engineering materials for structural applications. However, inhomogeneous deformation and severe shear localization at ambient temperature often lead to early failure and limit their reliability as structural materials. Heat treatments around the glass transition temperature (Tg ) generally aggravate the intrinsic brittleness of BMGs. In this paper, we report on the evolution of a nanoscale inhomogeneous microstructure upon low-temperature annealing in a Zr55 Cu30 Al10 Ni5 BMG. This important outcome is explained by the experimentally observed tendency for chemical decomposition between Cu and Zr of the investigated amorphous system and is in accordance with literature data on various Zr–Cu-based amorphous alloys. Finally, these local fluctuations influence the plasticity of BMGs beneficially, in contrast of the generally accepted embrittlement upon annealing. © 2008 Elsevier B.V. All rights reserved. Keywords: Metallic glass; Heat treatment; Microstructure; Plasticity
1. Introduction Due to absence of long range atomic order in bulk metallic glasses (BMGs) offers them unique physical, chemical and mechanical properties, compared to conventional crystalline metallic materials, making them a promising class of engineering materials. Their exceptionally high yield strength and yield strain in particular offers them high potential for structural applications [1–4]. However, localized and inhomogeneous deformation at ambient temperature [5], accompanied with strain [6–8] and thermal softening [9] lead to early failure after a small amount of macroscopic deformation which still limits the reliability of BMGs for structural applications. To circumvent this catastrophic failure, recent investigation aims to develop ‘heterogeneous’ microstructures in different length scales. Both structural and chemical inhomogeneities, with length scales from micrometer-sized ductile dendrites to
∗
Corresponding author. E-mail address: [email protected] (N. Van Steenberge).
0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.01.083
atomic-scale inhomogeneities have proven their value in increasing plasticity as compared to the monolithic glass [10–16]. Presence of such inhomogeneities promotes the plastic deformability because they increase nucleation of shear bands, induce their branching, and/or hinder their rapid propagation due to delocalization of the shear bands [17–19]. The way to obtain these inhomogeneities can be done ex situ (by physically adding a second phase to the molten alloy) [10] or by in situ methods. The latter involves obtaining them directly upon casting by designing a proper alloy composition [11,12,14–16] or upon annealing above crystallization temperature [13] or at least within the supercooled liquid region (SLR) (above the glass transition temperature Tg ) [19]. Annealing of metallic glasses to temperatures below their Tg leads to structural relaxation, a process which is generally believed to be mainly characterized by annihilation of excess free volume [7,8,20], leading to a deterioration of plasticity [21–23]. Very recently however, subtle structural changes by low-temperature annealing (
N. Van Steenberge et al. / Materials Science and Engineering A 491 (2008) 124–130
In this paper, we report on the synthesis of an inhomogeneous microstructure in Zr55 Cu30 Al10 Ni5 BMG upon annealing below Tg . We observed that the evolution of clusters through both macroscopic and microscopic investigation by synchrotron X-ray diffraction and differential scanning calorimetry (DSC), and by transmission electron microscopy (TEM), respectively, both ex situ (after annealing), combined with high-angle annular dark field scanning transmission electron microscopy (HAADFSTEM), as well as during in situ heat treatments in the TEM. Furthermore, uniaxial compression tests show that plasticity increases significantly in the annealed sample—in contrast with the generally accepted deterioration of plasticity [22,23].
2. Experimental conditions Rods of 3 mm diameter and 50 mm length were obtained by injection casting of arc-melted Zr55 Cu30 Al10 Ni5 (at.%) ingots into a Cu-mould. Heat treatments on several rods were performed under vacuum with Ti as oxygen-getter in a tubular furnace, introducing the samples into a preheated furnace at 623 K, which is below the glass transition. Once the samples reached the desired temperature, they were immediately cooled down to room temperature. Oxygen content was determined both after casting (before annealing) and after annealing by carrier gas-hot extraction with a O/N analyzer TC-436/LECO USA. The amorphous nature of all samples was verified by X-ray diffraction (XRD), using a Philips X’Pert diffractometer with monochromatic Cu K␣ radiation. To analyze the shortand medium-range order, high-energy X-ray diffraction experiments were conducted at the synchrotron beam-line BW 5 (HASYLAB, Hamburg), using a wavelength of λ = 0.0125 nm. For thermal analysis, a PerkinElmer model DSC-7 was used at a heating rate of 0.667 K/s. A 200 kV Hitachi 600AB TEM was used for microstructural characterization of as-cast and heat treated samples while a 300 kV Tecnai F30 with field emission was used for high-angle annular dark field (HAADF) imaging and energy dispersive spectroscopic compositional analysis (EDX) in scanning transmission mode (STEM). In situ TEM annealing experiments were performed in a JEOL JEM-2011 at a rather low heating rate of 0.333 K/s (to avoid overshoot) up to 623 K and holding segment of 30 min at the same temperature. The sample holder, designed by Gatan, has a special design to minimize thermal drift during the experiment. The experiment was repeated with a defocused beam at low magnification during the heating part to minimize the introduction of beam damage in the sample, comparing the initial image (before heating) with the images made at 623 K. Afterwards, the samples were further analyzed by electron energy loss spectroscopy (EELS), using a 200 kV F20-SACTEM Tecnai TEM fitted with an energy filter GIF-Tridiem. Samples for TEM were mechanically thinned to 30 m and then further thinned by electropolishing them using a solution of 30% HNO3 in methanol at a temperature of 253 K. Due caution was taken to maintain the same conditions for the preparation of each sample. Afterwards, ion beam thinning at 2 keV and an angle of 4◦ was applied for a short time.
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Uni-axial compression tests were carried out using an Instron 8562 testing machine at room temperature. The compression test samples were prepared according to ASTM standards to a perfectly orthogonal geometry with an aspect ratio of 2:1. In total up to 20 compression tests were performed with at least four rods for each condition, taken from bottom and centre parts along the rod. Nanoindentation experiments were carried out with an MTS nanoindenter XP, at room temperature, in load control mode using a Berkovich indenter tip. Prior to nanoindentation, the samples were carefully polished to mirror-like appearance using diamond paste. The indentation function consisted of a loading segment of 40 s, followed by a load holding segment of 30 s and an unloading segment of 10 s. The maximum applied load was 500 mN. A Jeol JSM-6300 SEM was used to observe shear bands on the outer surface of the compressed samples and around the nanoindents. 3. Results and discussion Investigation by conventional XRD reveals the ‘Bragg amorphous’ nature of both the as-cast as the annealed samples: two broad diffuse halos are observed, without any sharp diffraction peaks typical for the presence of crystalline phases (Fig. 1a). Fitting the main halos to a Lorentzian function does not reveal any difference in halo widths. Calculation of the pair correlation function g(r) (Fig. 1b–d) based on high-energy X-ray diffraction measurements, as described, e.g. in Ref. [25], however, shows small but significant changes in the peak profile of the annealed sample compared to the as-cast sample. This figure also contains data on an annealed and deformed sample which will be discussed below. The main peaks at interatomic distances of 0.275 and 0.312 nm can be attributed to the pair between Cu–Zr and Zr–Zr, respectively, based on literature data for binary Zr–Cubased alloys [25]. Upon annealing, the intensity of Cu–Zr pair decreases, while the peak for Zr–Zr pair intensifies. Furthermore, changes in the second peak around 0.5 nm indicate a change in medium range order in the glassy phase. The initially flat peak for the as-cast sample starts to split after annealing. Based on the partial pair correlation functions for Cu–Cu, Zr–Cu and Zr–Zr interactions made by Mattern et al. [25], the intensity of Zr–Cupairs decreases, while there is a slight increase in probability for the other two pairs to occur. These observations are in good agreement with the observations after low-temperature annealing of a Zr-based BMG, as made by neutron diffraction [24], although the changes observed in this work occurred on an even smaller time-scale. Thermal analysis confirms the general preservation of the amorphous nature upon annealing (Fig. 2). The transition from glass to supercooled liquid occurs at 702 ± 4 K while the width of the supercooled liquid region is 79 ± 2 K, for a heating rate of 0.667 K/s. Comparison of the top and the bottom parts of the as-cast rods revealed very good homogeneity along the length of the rod. Upon annealing, the samples relaxed to a more thermodynamically stable state, resulting in a lower enthalpy release during the exothermic event preceding the glass transition. By integration of the crystallization peak area, values of −52 ± 2 and −54 ± 2 J/g calculated respectively for the as-cast sample
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Fig. 1. Effect of the heat treatment on the short–medium range order. (a) The conventional X-ray diffraction spectrum of the as-cast sample compared to the spectra of the annealed sample shows the typical “Bragg amorphous” structure. (b) Pair correlation function g(r) = ρ(r)/ρ0 of the as-cast and the annealed sample, together with the pair correlation function of an annealed and deformed sample. Details of the first (c) and second (d) “peak” in (b) show small but significant changes in short–medium range order.
and the annealed sample. On the other hand, an additional shoulder appears in the crystallization peak of the annealed sample (inset of Fig. 2). This change in shape has been attributed in literature to the occurrence of phase separation preceding the crystallization process [26], formation of clusters upon a preceding deformation treatment [27] as well as to the presence of oxygen [28]. However, the latter possibility can be ruled out since the oxygen content was measured to be around 0.0005 at.% both before and after annealing treatment. In this work, it is related to the formation of clusters as well, as shown below.
Fig. 2. Effect of the annealing on thermal stability and crystallization behaviour. The DSC trace of the as-cast specimen is shown for comparison. After annealing below Tg , splitting of the crystallization peak is observed (inset).
Imaging by TEM reveals a clear difference in the microstructures of the as-cast and annealed sample (Fig. 3). While the as-cast sample (a) shows a homogeneous, featureless microstructure, the microstructure of annealed specimen (b) consists of a bright matrix with darker regions on the order of 5 nm. Their respective selected area electron diffraction (SAED) patterns on the other hand do not show a significant change. The contrast was observed all over the sample. A small amount of dark zones with an average size of 20–30 nm is also present. In some of these darker regions finally, lattice fringes in the range of 2–5 nm, pointing towards the presence of nanocrystallites, were detected by high-resolution TEM. Preparation of samples for TEM sometimes leads to inhomogeneous thinning, which might be interpreted erroneously as compositional contrast [29]. To rule out this possibility, the sample was checked by HAADF, where contrast is based mainly on the atomic number, and depends less on thickness. Furthermore, annealing experiments in situ during TEM observation with artefact-free samples can help to gain insight into microstructural changes upon heating. HAADF-STEM imaging of the annealed sample confirms that the observed contrast in Fig. 3b is due to compositional fluctuations (Fig. 4a) [16,30] and is not the result of an inhomogeneous thickness. Mapping by EDX in scanning transmission mode reveals that the dark zones in Fig. 3b (bright zones in STEM image) are enriched in Cu (Fig. 4b and c), while no distinct variations of the other elements were observed (Fig. 4d–f). In situ annealing in the TEM shows the evolution of an inhomogeneous microstructure with brighter regions and darker borders upon annealing from an initially homogeneous glassy
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Fig. 3. Microstructural features in the as-cast (a) and annealed (623 K) sample (b), observed by TEM, revealing the change of a homogeneous microstructure to a heterogeneous one upon annealing. The respective SAED patterns are shown in the insets.
microstructure, as shown in Fig. 5a (as-cast) and b (623 K). The SAED spectra of the respective samples show clearly a change in the diffraction pattern: the initially broad diffuse halo changes into two sharper rings, pointing towards the development of a more ordered structure upon annealing. EELS measurements in these in situ annealed samples is in agreement with the EDX-STEM results, showing an enrichment of Cu in the darker zones (not shown here). In the darker ‘edges’ nanocrystals started to appear after annealing for 20 min, which could be observed in dark field (not shown here). The brighter
regions were stable and no crystallization occurred during the course of this experiment (30 min) below Tg . Annealing to crystallization temperature shows the formation of these white areas, which initially are amorphous, but crystallize within a few minutes at this higher temperature. Since it is not possible to observe a large area during the recording of the in situ experiments, the first image of such inhomogeneous microstructure was observed after holding for 10 min and was located above the initial region of observation. However, based on the size of the bright regions, already arisen in such a short period of time, it is most likely
Fig. 4. High-angle annular dark field image (HAADF) shows an overall compositional contrast (a). Mapping by EDX is made in detail in one of the bright areas (b), indicated by the square in (a), of the respective elements (c–f), a clear local enrichment in Cu and (c) is observed in the brighter regions of (a), while the distribution of other elements (Zr, Ni and Al), is rather homogeneous (d–f).
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Fig. 5. Microstructural features in the as-cast (a) and in the annealed sample (b), observed in situ during annealing in the TEM. These results show the tendency of this glass to decompose upon annealing. The SAED patterns show the gradual creation of a more ordered structure upon annealing.
to assume that chemical fluctuations occur earlier on a smaller scale, which might have been too small to be detected with the magnification and microscope set-up used for these experiments. The authors are also aware that diffusion during these in situ experiments was probably enhanced by surface defects which might also explain the large size of the bright areas observed. Nonetheless, the important information obtained here is that this system shows a strong tendency to decompose as first step towards structural ordering, even already below Tg . This observation is in accordance with the reduced probability for Zr–Cu interactions observed by synchrotron radiation X-ray diffraction experiments. In fact, systems based on Zr–Cu have been reported several times in literature to show a strong tendency for decomposition into Cu- and Zr-rich regions, although it might not be expected based on the available data of their heat of mixing [15]. For example, such decomposition has been reported in the binary system over the whole composition range of Cu-25–65 at.% Zr directly during synthesis when the sample was prepared by vapour-quenching and after annealing when the sample was prepared by quenching from the liquid state [31]. Decomposition and formation of Cu-rich clusters in similar compositions have further been observed upon (high temperature) deformation [27] and after hydrogenation experiments [32]. This leads to the conclusion that this decomposed state must be more thermodynamically stable than the monolithic glass, independent whether the alloy is in its glassy state or is a (supercooled) liquid. Indeed, experiments with Zr–Al–Cu–Ni diffusion couples combined with thermodynamic calculations [33] showed that
Cu has no tendency at all to mix with Zr(Al, Ni) upon annealing. On the other hand, the diffusion constants of Ni (and Cu) seemed to be mainly influenced by the thermodynamic driving force to lower the system’s Gibbs free energy, rather than by the fact whether the alloy was above or below the glass transition. Although the structural changes occurring upon annealing are moderate, they turn out to have a rather important influence on the plasticity in compression. Fig. 6a shows the true stress–strain curve after uni-axial compression of the as-cast and heat-treated sample, annealed to 623 K. In both cases, an average yield stress of around 1805 ± 52 MPa and an elastic strain of 2.1% were obtained. However, unlike the reported embrittlement upon annealing [21–23], an increase of the total strain from 2.7 ± 0.5% for the as-cast sample up to 10 ± 3% for the annealed samples is observed. Furthermore, the annealed sample shows a weak increase of the flow stress [34]. Improvement of the plasticity must have been provoked by a multiplication of shear bands in the glassy phase. Indeed, the outer surface of the compressed annealed sample shows a large number of shear bands, whereas the as-cast samples showed only one or two principal shear bands that are crossing over the specimen (SEM images given in the inset of Fig. 6a). Moreover, strong branching, deflection and interaction took place between the shear bands developed in the annealed sample. Similarly, upon nanoindentation, a larger amount of shear bands surrounding and inside the indent are observed in the annealed sample as compared to the as-cast sample (Fig. 6b). Multiplication of shear bands during nanoindentation has been reported a few times in literature under various conditions and different explanations were given to it. Li et al [35] observed a similar multiplication of shear bands during nanoindentation at cryogenic temperature. They concluded that due to the reduced propagation possibilities of the shear bands, such a multiplication was necessary in order to accommodate the applied strain during nanoindentation. Jiang and Atzmon [36] attributed the increased amount of shear bands to a required increase in nucleation rate, due to a reduction of free volume upon annealing. Besides reduction of free volume, we believe that the presence of structural changes upon annealing might have contributed to the observed multiplication. Moreover, Dmowski et al. showed that the first step during structural relaxation cannot be correlated to annihilation of free volume, but should involve subtle structural changes [24]. The evolution of compositional fluctuations upon annealing under the form of Cu-clusters must results in local fluctuations in elastic properties. The cooperative shear model for yielding of metallic glasses [37] foresees that the barrier height for shear flow for a given glass is proportional to the shear modulus G. Furthermore, recent simulations predict a decrease of G with decreasing Cu-content in Cu–Zr-alloys [38]. The Zr-based matrix, which is depleted to a certain extent with Cu, is less stiff and will therefore carry the plastic deformation, i.e. shear bands can easier nucleate. The Cu-rich clusters on the other hand show probably a higher degree of order (more solid-like) and are stiffer, therefore forming a barrier for shear deformation. On the other hand, further structural changes occur during deformation, as can be deduced from Fig. 1d. Therefore, the exact influence
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Fig. 6. (a) Comparison of plastic flow behaviours in uni-axial compression of an as-cast sample and annealed sample (
of the microstructural changes observed upon annealing on the increase of plasticity is still under investigation. 4. Conclusions In summary, both synchrotron radiation diffraction experiments and in situ annealing experiments during TEM observation, show the tendency for decomposition between Cu and Zr of the investigated ZrCuAlNi amorphous system, in accordance with literature data on various Zr–Cu-based amorphous alloys. This tendency for chemical decomposition even leads to the evolution of nanoscale inhomogeneities (Cuclusters) upon short-time annealing below the glass transition temperature. These local compositional fluctuations have a rather important influence on the improvement of plasticity of BMGs, without a significant loss of strength, in contrast of the generally accepted embrittlement upon annealing.
Acknowledgements The authors thank the technical support by S. Kuzinski, M. Gr¨undlich, H.-J. Klau and J. Thomas at IFW Dresden, J. Portillo from the Servei Cientificot`ecnic from the Universitat de Barcelona as well as E. Rossinyol and O. Castells from the Servei de Microsc`opia from the Universitat Aut`onoma de Barcelona. Furthermore, we are grateful for the technical support by F. Mompiou during the in situ TEM experiments, performed at the CEMES in Toulouse with the financial support of the European Commission within the framework of the ESTEEM (Enabling Science and Technology through European Electron Microscopy) project (026019). NVS acknowledges further the financial support from the European Commission under MCRTN contract “Ductile bulk metallic glass composites” (MRTN-CT-2003-504692) and the special support action “Strengthening the role of women scientists in nano-science”.
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Partial financial support from UAB (PNL 2006-019), DURSI (2005-SGR-00401) and MEC (MAT 2007-61629) projects is also acknowledged. References [1] A. Peker, W.L. Johnson, Appl. Phys. Lett. 63 (1993) 2342–2344. [2] A. Inoue, T. Zhang, J. Non-Cryst. Solids 250 (1999) 552–559. [3] A. Inoue, B. Shen, H. Koshiba, H. Kato, A.R. Yavari, Nat. Mater. 2 (2003) 661–663. [4] M.F. Ashby, A.L. Greer, Scripta Mater. 54 (2006) 321–326. [5] F. Spaepen, Acta Metall. 25 (1977) 407–415. [6] H. Bei, S. Xie, E.P. George, Phys. Rev. Lett. 96 (2006) 105503. [7] B.P. Kanungo, S.C. Glade, P. Asoka-Kumar, K.M. Flores, Intermetallics 12 (2004) 1073–1080. [8] K. Hajlaoui, T. Benameur, G. Vaughan, A.R. Yavari, Scripta Mater. 51 (2004) 843–848. [9] J.J. Lewandowski, A.L. Greer, Nat. Mater. 5 (2006) 15–18. [10] H. Choi-Yim, R.D. Conner, F. Szuecs, W.L. Johnson, Acta Mater. 50 (2002) 2737–2745. [11] U. Kuhn, J. Eckert, N. Mattern, L. Schultz, Mat. Sci. Eng. A 375–377 (2004) 322–326. [12] H. Kato, T. Hirano, A. Matsuo, Y. Kawamura, A. Inoue, Scripta Mater. 43 (2000) 503–507. [13] C. Fan, A. Inoue, Appl. Phys. Lett. 77 (2000) 46–48. [14] J.C. Oh, T. Ohkubo, Y.C. Kim, E. Fleury, K. Hono, Scripta Mater. 53 (2005) 165–169. [15] E.S. Park, D.H. Kim, Acta Mater. 54 (2006) 2597–2604. [16] K.B. Kim, J. Das, F. Baier, M.B. Tang, W.H. Wang, J. Eckert, Appl. Phys. Lett. 88 (2006) 051911. [17] R.T. Ott, F. Sansoz, J.F. Molinari, K.T. Ramesh, T.C. Hufnagel, Acta Mater. 53 (2005) 1883–1893. [18] A. Concustell, N. Mattern, H. Wendrock, U. Kuehn, A. Gebert, J. Eckert, A.L. Greer, J. Sort, M.D. Bar´o, Scripta Mater. 56 (2007) 85–88.
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Microstructural inhomogeneities introduced in a Zr-based bulk metallic glass upon low-temperature annealing N. Van Steenberge a,∗ , A. Concustell a , J. Sort b , J. Das c , N. Mattern c , A. Gebert c , S. Suri˜nach a , J. Eckert c , M.D. Bar´o a b
a Departament de F´ısica, Universitat Aut` onoma de Barcelona, 08913 Bellaterra, Spain Instituci´o Catalana de Recerca i Estudis Avan¸cats and Departament de F´ısica, Universitat Aut`onoma de Barcelona, 08193 Bellaterra, Spain c Leibniz-Institut f¨ ur Festk¨orper- und Werkstoffforschung Dresden, P.O. Box 27 00 16, 01171 Dresden, Germany
Received 19 October 2007; received in revised form 21 January 2008; accepted 25 January 2008
Abstract Due to their exceptionally high yield strength and yield strain as compared to conventional metallic materials, bulk metallic glasses (BMGs) represent a class of promising engineering materials for structural applications. However, inhomogeneous deformation and severe shear localization at ambient temperature often lead to early failure and limit their reliability as structural materials. Heat treatments around the glass transition temperature (Tg ) generally aggravate the intrinsic brittleness of BMGs. In this paper, we report on the evolution of a nanoscale inhomogeneous microstructure upon low-temperature annealing in a Zr55 Cu30 Al10 Ni5 BMG. This important outcome is explained by the experimentally observed tendency for chemical decomposition between Cu and Zr of the investigated amorphous system and is in accordance with literature data on various Zr–Cu-based amorphous alloys. Finally, these local fluctuations influence the plasticity of BMGs beneficially, in contrast of the generally accepted embrittlement upon annealing. © 2008 Elsevier B.V. All rights reserved. Keywords: Metallic glass; Heat treatment; Microstructure; Plasticity
1. Introduction Due to absence of long range atomic order in bulk metallic glasses (BMGs) offers them unique physical, chemical and mechanical properties, compared to conventional crystalline metallic materials, making them a promising class of engineering materials. Their exceptionally high yield strength and yield strain in particular offers them high potential for structural applications [1–4]. However, localized and inhomogeneous deformation at ambient temperature [5], accompanied with strain [6–8] and thermal softening [9] lead to early failure after a small amount of macroscopic deformation which still limits the reliability of BMGs for structural applications. To circumvent this catastrophic failure, recent investigation aims to develop ‘heterogeneous’ microstructures in different length scales. Both structural and chemical inhomogeneities, with length scales from micrometer-sized ductile dendrites to
∗
Corresponding author. E-mail address: [email protected] (N. Van Steenberge).
0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.01.083
atomic-scale inhomogeneities have proven their value in increasing plasticity as compared to the monolithic glass [10–16]. Presence of such inhomogeneities promotes the plastic deformability because they increase nucleation of shear bands, induce their branching, and/or hinder their rapid propagation due to delocalization of the shear bands [17–19]. The way to obtain these inhomogeneities can be done ex situ (by physically adding a second phase to the molten alloy) [10] or by in situ methods. The latter involves obtaining them directly upon casting by designing a proper alloy composition [11,12,14–16] or upon annealing above crystallization temperature [13] or at least within the supercooled liquid region (SLR) (above the glass transition temperature Tg ) [19]. Annealing of metallic glasses to temperatures below their Tg leads to structural relaxation, a process which is generally believed to be mainly characterized by annihilation of excess free volume [7,8,20], leading to a deterioration of plasticity [21–23]. Very recently however, subtle structural changes by low-temperature annealing (
N. Van Steenberge et al. / Materials Science and Engineering A 491 (2008) 124–130
In this paper, we report on the synthesis of an inhomogeneous microstructure in Zr55 Cu30 Al10 Ni5 BMG upon annealing below Tg . We observed that the evolution of clusters through both macroscopic and microscopic investigation by synchrotron X-ray diffraction and differential scanning calorimetry (DSC), and by transmission electron microscopy (TEM), respectively, both ex situ (after annealing), combined with high-angle annular dark field scanning transmission electron microscopy (HAADFSTEM), as well as during in situ heat treatments in the TEM. Furthermore, uniaxial compression tests show that plasticity increases significantly in the annealed sample—in contrast with the generally accepted deterioration of plasticity [22,23].
2. Experimental conditions Rods of 3 mm diameter and 50 mm length were obtained by injection casting of arc-melted Zr55 Cu30 Al10 Ni5 (at.%) ingots into a Cu-mould. Heat treatments on several rods were performed under vacuum with Ti as oxygen-getter in a tubular furnace, introducing the samples into a preheated furnace at 623 K, which is below the glass transition. Once the samples reached the desired temperature, they were immediately cooled down to room temperature. Oxygen content was determined both after casting (before annealing) and after annealing by carrier gas-hot extraction with a O/N analyzer TC-436/LECO USA. The amorphous nature of all samples was verified by X-ray diffraction (XRD), using a Philips X’Pert diffractometer with monochromatic Cu K␣ radiation. To analyze the shortand medium-range order, high-energy X-ray diffraction experiments were conducted at the synchrotron beam-line BW 5 (HASYLAB, Hamburg), using a wavelength of λ = 0.0125 nm. For thermal analysis, a PerkinElmer model DSC-7 was used at a heating rate of 0.667 K/s. A 200 kV Hitachi 600AB TEM was used for microstructural characterization of as-cast and heat treated samples while a 300 kV Tecnai F30 with field emission was used for high-angle annular dark field (HAADF) imaging and energy dispersive spectroscopic compositional analysis (EDX) in scanning transmission mode (STEM). In situ TEM annealing experiments were performed in a JEOL JEM-2011 at a rather low heating rate of 0.333 K/s (to avoid overshoot) up to 623 K and holding segment of 30 min at the same temperature. The sample holder, designed by Gatan, has a special design to minimize thermal drift during the experiment. The experiment was repeated with a defocused beam at low magnification during the heating part to minimize the introduction of beam damage in the sample, comparing the initial image (before heating) with the images made at 623 K. Afterwards, the samples were further analyzed by electron energy loss spectroscopy (EELS), using a 200 kV F20-SACTEM Tecnai TEM fitted with an energy filter GIF-Tridiem. Samples for TEM were mechanically thinned to 30 m and then further thinned by electropolishing them using a solution of 30% HNO3 in methanol at a temperature of 253 K. Due caution was taken to maintain the same conditions for the preparation of each sample. Afterwards, ion beam thinning at 2 keV and an angle of 4◦ was applied for a short time.
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Uni-axial compression tests were carried out using an Instron 8562 testing machine at room temperature. The compression test samples were prepared according to ASTM standards to a perfectly orthogonal geometry with an aspect ratio of 2:1. In total up to 20 compression tests were performed with at least four rods for each condition, taken from bottom and centre parts along the rod. Nanoindentation experiments were carried out with an MTS nanoindenter XP, at room temperature, in load control mode using a Berkovich indenter tip. Prior to nanoindentation, the samples were carefully polished to mirror-like appearance using diamond paste. The indentation function consisted of a loading segment of 40 s, followed by a load holding segment of 30 s and an unloading segment of 10 s. The maximum applied load was 500 mN. A Jeol JSM-6300 SEM was used to observe shear bands on the outer surface of the compressed samples and around the nanoindents. 3. Results and discussion Investigation by conventional XRD reveals the ‘Bragg amorphous’ nature of both the as-cast as the annealed samples: two broad diffuse halos are observed, without any sharp diffraction peaks typical for the presence of crystalline phases (Fig. 1a). Fitting the main halos to a Lorentzian function does not reveal any difference in halo widths. Calculation of the pair correlation function g(r) (Fig. 1b–d) based on high-energy X-ray diffraction measurements, as described, e.g. in Ref. [25], however, shows small but significant changes in the peak profile of the annealed sample compared to the as-cast sample. This figure also contains data on an annealed and deformed sample which will be discussed below. The main peaks at interatomic distances of 0.275 and 0.312 nm can be attributed to the pair between Cu–Zr and Zr–Zr, respectively, based on literature data for binary Zr–Cubased alloys [25]. Upon annealing, the intensity of Cu–Zr pair decreases, while the peak for Zr–Zr pair intensifies. Furthermore, changes in the second peak around 0.5 nm indicate a change in medium range order in the glassy phase. The initially flat peak for the as-cast sample starts to split after annealing. Based on the partial pair correlation functions for Cu–Cu, Zr–Cu and Zr–Zr interactions made by Mattern et al. [25], the intensity of Zr–Cupairs decreases, while there is a slight increase in probability for the other two pairs to occur. These observations are in good agreement with the observations after low-temperature annealing of a Zr-based BMG, as made by neutron diffraction [24], although the changes observed in this work occurred on an even smaller time-scale. Thermal analysis confirms the general preservation of the amorphous nature upon annealing (Fig. 2). The transition from glass to supercooled liquid occurs at 702 ± 4 K while the width of the supercooled liquid region is 79 ± 2 K, for a heating rate of 0.667 K/s. Comparison of the top and the bottom parts of the as-cast rods revealed very good homogeneity along the length of the rod. Upon annealing, the samples relaxed to a more thermodynamically stable state, resulting in a lower enthalpy release during the exothermic event preceding the glass transition. By integration of the crystallization peak area, values of −52 ± 2 and −54 ± 2 J/g calculated respectively for the as-cast sample
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Fig. 1. Effect of the heat treatment on the short–medium range order. (a) The conventional X-ray diffraction spectrum of the as-cast sample compared to the spectra of the annealed sample shows the typical “Bragg amorphous” structure. (b) Pair correlation function g(r) = ρ(r)/ρ0 of the as-cast and the annealed sample, together with the pair correlation function of an annealed and deformed sample. Details of the first (c) and second (d) “peak” in (b) show small but significant changes in short–medium range order.
and the annealed sample. On the other hand, an additional shoulder appears in the crystallization peak of the annealed sample (inset of Fig. 2). This change in shape has been attributed in literature to the occurrence of phase separation preceding the crystallization process [26], formation of clusters upon a preceding deformation treatment [27] as well as to the presence of oxygen [28]. However, the latter possibility can be ruled out since the oxygen content was measured to be around 0.0005 at.% both before and after annealing treatment. In this work, it is related to the formation of clusters as well, as shown below.
Fig. 2. Effect of the annealing on thermal stability and crystallization behaviour. The DSC trace of the as-cast specimen is shown for comparison. After annealing below Tg , splitting of the crystallization peak is observed (inset).
Imaging by TEM reveals a clear difference in the microstructures of the as-cast and annealed sample (Fig. 3). While the as-cast sample (a) shows a homogeneous, featureless microstructure, the microstructure of annealed specimen (b) consists of a bright matrix with darker regions on the order of 5 nm. Their respective selected area electron diffraction (SAED) patterns on the other hand do not show a significant change. The contrast was observed all over the sample. A small amount of dark zones with an average size of 20–30 nm is also present. In some of these darker regions finally, lattice fringes in the range of 2–5 nm, pointing towards the presence of nanocrystallites, were detected by high-resolution TEM. Preparation of samples for TEM sometimes leads to inhomogeneous thinning, which might be interpreted erroneously as compositional contrast [29]. To rule out this possibility, the sample was checked by HAADF, where contrast is based mainly on the atomic number, and depends less on thickness. Furthermore, annealing experiments in situ during TEM observation with artefact-free samples can help to gain insight into microstructural changes upon heating. HAADF-STEM imaging of the annealed sample confirms that the observed contrast in Fig. 3b is due to compositional fluctuations (Fig. 4a) [16,30] and is not the result of an inhomogeneous thickness. Mapping by EDX in scanning transmission mode reveals that the dark zones in Fig. 3b (bright zones in STEM image) are enriched in Cu (Fig. 4b and c), while no distinct variations of the other elements were observed (Fig. 4d–f). In situ annealing in the TEM shows the evolution of an inhomogeneous microstructure with brighter regions and darker borders upon annealing from an initially homogeneous glassy
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Fig. 3. Microstructural features in the as-cast (a) and annealed (623 K) sample (b), observed by TEM, revealing the change of a homogeneous microstructure to a heterogeneous one upon annealing. The respective SAED patterns are shown in the insets.
microstructure, as shown in Fig. 5a (as-cast) and b (623 K). The SAED spectra of the respective samples show clearly a change in the diffraction pattern: the initially broad diffuse halo changes into two sharper rings, pointing towards the development of a more ordered structure upon annealing. EELS measurements in these in situ annealed samples is in agreement with the EDX-STEM results, showing an enrichment of Cu in the darker zones (not shown here). In the darker ‘edges’ nanocrystals started to appear after annealing for 20 min, which could be observed in dark field (not shown here). The brighter
regions were stable and no crystallization occurred during the course of this experiment (30 min) below Tg . Annealing to crystallization temperature shows the formation of these white areas, which initially are amorphous, but crystallize within a few minutes at this higher temperature. Since it is not possible to observe a large area during the recording of the in situ experiments, the first image of such inhomogeneous microstructure was observed after holding for 10 min and was located above the initial region of observation. However, based on the size of the bright regions, already arisen in such a short period of time, it is most likely
Fig. 4. High-angle annular dark field image (HAADF) shows an overall compositional contrast (a). Mapping by EDX is made in detail in one of the bright areas (b), indicated by the square in (a), of the respective elements (c–f), a clear local enrichment in Cu and (c) is observed in the brighter regions of (a), while the distribution of other elements (Zr, Ni and Al), is rather homogeneous (d–f).
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Fig. 5. Microstructural features in the as-cast (a) and in the annealed sample (b), observed in situ during annealing in the TEM. These results show the tendency of this glass to decompose upon annealing. The SAED patterns show the gradual creation of a more ordered structure upon annealing.
to assume that chemical fluctuations occur earlier on a smaller scale, which might have been too small to be detected with the magnification and microscope set-up used for these experiments. The authors are also aware that diffusion during these in situ experiments was probably enhanced by surface defects which might also explain the large size of the bright areas observed. Nonetheless, the important information obtained here is that this system shows a strong tendency to decompose as first step towards structural ordering, even already below Tg . This observation is in accordance with the reduced probability for Zr–Cu interactions observed by synchrotron radiation X-ray diffraction experiments. In fact, systems based on Zr–Cu have been reported several times in literature to show a strong tendency for decomposition into Cu- and Zr-rich regions, although it might not be expected based on the available data of their heat of mixing [15]. For example, such decomposition has been reported in the binary system over the whole composition range of Cu-25–65 at.% Zr directly during synthesis when the sample was prepared by vapour-quenching and after annealing when the sample was prepared by quenching from the liquid state [31]. Decomposition and formation of Cu-rich clusters in similar compositions have further been observed upon (high temperature) deformation [27] and after hydrogenation experiments [32]. This leads to the conclusion that this decomposed state must be more thermodynamically stable than the monolithic glass, independent whether the alloy is in its glassy state or is a (supercooled) liquid. Indeed, experiments with Zr–Al–Cu–Ni diffusion couples combined with thermodynamic calculations [33] showed that
Cu has no tendency at all to mix with Zr(Al, Ni) upon annealing. On the other hand, the diffusion constants of Ni (and Cu) seemed to be mainly influenced by the thermodynamic driving force to lower the system’s Gibbs free energy, rather than by the fact whether the alloy was above or below the glass transition. Although the structural changes occurring upon annealing are moderate, they turn out to have a rather important influence on the plasticity in compression. Fig. 6a shows the true stress–strain curve after uni-axial compression of the as-cast and heat-treated sample, annealed to 623 K. In both cases, an average yield stress of around 1805 ± 52 MPa and an elastic strain of 2.1% were obtained. However, unlike the reported embrittlement upon annealing [21–23], an increase of the total strain from 2.7 ± 0.5% for the as-cast sample up to 10 ± 3% for the annealed samples is observed. Furthermore, the annealed sample shows a weak increase of the flow stress [34]. Improvement of the plasticity must have been provoked by a multiplication of shear bands in the glassy phase. Indeed, the outer surface of the compressed annealed sample shows a large number of shear bands, whereas the as-cast samples showed only one or two principal shear bands that are crossing over the specimen (SEM images given in the inset of Fig. 6a). Moreover, strong branching, deflection and interaction took place between the shear bands developed in the annealed sample. Similarly, upon nanoindentation, a larger amount of shear bands surrounding and inside the indent are observed in the annealed sample as compared to the as-cast sample (Fig. 6b). Multiplication of shear bands during nanoindentation has been reported a few times in literature under various conditions and different explanations were given to it. Li et al [35] observed a similar multiplication of shear bands during nanoindentation at cryogenic temperature. They concluded that due to the reduced propagation possibilities of the shear bands, such a multiplication was necessary in order to accommodate the applied strain during nanoindentation. Jiang and Atzmon [36] attributed the increased amount of shear bands to a required increase in nucleation rate, due to a reduction of free volume upon annealing. Besides reduction of free volume, we believe that the presence of structural changes upon annealing might have contributed to the observed multiplication. Moreover, Dmowski et al. showed that the first step during structural relaxation cannot be correlated to annihilation of free volume, but should involve subtle structural changes [24]. The evolution of compositional fluctuations upon annealing under the form of Cu-clusters must results in local fluctuations in elastic properties. The cooperative shear model for yielding of metallic glasses [37] foresees that the barrier height for shear flow for a given glass is proportional to the shear modulus G. Furthermore, recent simulations predict a decrease of G with decreasing Cu-content in Cu–Zr-alloys [38]. The Zr-based matrix, which is depleted to a certain extent with Cu, is less stiff and will therefore carry the plastic deformation, i.e. shear bands can easier nucleate. The Cu-rich clusters on the other hand show probably a higher degree of order (more solid-like) and are stiffer, therefore forming a barrier for shear deformation. On the other hand, further structural changes occur during deformation, as can be deduced from Fig. 1d. Therefore, the exact influence
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Fig. 6. (a) Comparison of plastic flow behaviours in uni-axial compression of an as-cast sample and annealed sample (
of the microstructural changes observed upon annealing on the increase of plasticity is still under investigation. 4. Conclusions In summary, both synchrotron radiation diffraction experiments and in situ annealing experiments during TEM observation, show the tendency for decomposition between Cu and Zr of the investigated ZrCuAlNi amorphous system, in accordance with literature data on various Zr–Cu-based amorphous alloys. This tendency for chemical decomposition even leads to the evolution of nanoscale inhomogeneities (Cuclusters) upon short-time annealing below the glass transition temperature. These local compositional fluctuations have a rather important influence on the improvement of plasticity of BMGs, without a significant loss of strength, in contrast of the generally accepted embrittlement upon annealing.
Acknowledgements The authors thank the technical support by S. Kuzinski, M. Gr¨undlich, H.-J. Klau and J. Thomas at IFW Dresden, J. Portillo from the Servei Cientificot`ecnic from the Universitat de Barcelona as well as E. Rossinyol and O. Castells from the Servei de Microsc`opia from the Universitat Aut`onoma de Barcelona. Furthermore, we are grateful for the technical support by F. Mompiou during the in situ TEM experiments, performed at the CEMES in Toulouse with the financial support of the European Commission within the framework of the ESTEEM (Enabling Science and Technology through European Electron Microscopy) project (026019). NVS acknowledges further the financial support from the European Commission under MCRTN contract “Ductile bulk metallic glass composites” (MRTN-CT-2003-504692) and the special support action “Strengthening the role of women scientists in nano-science”.
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