- Email: [email protected]
Thermal embrittlement of Fe-based amorphous ribbons
Available online at www.sciencedirect.com
Journal of Non-Crystalline Solids 354 (2008) 882–888 www.elsevier.com/locate/jnoncrysol
Thermal embrittlement of Fe-based amorphous ribbons G. Kumar *, M. Ohnuma, T. Furubayashi, T. Ohkubo, K. Hono National Institute for Materials Science, 1-2-1 Sengen, Tsukuba 305-0047, Japan Received 6 April 2007; received in revised form 31 July 2007 Available online 14 September 2007
Abstract Thermal embrittlement of amorphous Fe78 xNixSi10B12 and Co78Si10B12 alloys has been studied. The majority of Fe-based amorphous alloys became brittle after annealing to temperatures significantly lower than their crystallization temperatures, whereas the non-Fe-based amorphous alloys remained ductile after annealing up to their crystallization temperatures. The present results emphasize on the key difference between the embrittlement behavior of Fe-based and non-Fe-based amorphous alloys. We point out a possible correlation between different phenomena like Invar effect, compositional inhomogenity, and residual stress playing a critical role in the thermal embrittlement of Fe-based amorphous alloys. 2007 Elsevier B.V. All rights reserved. PACS: 65.60.+a; 75.50.Bb; 81.05.Kf Keywords: Amorphous metals, metallic glasses; Transition metals; Mechanical, stress relaxation; Short-range order
1. Introduction Fe-based amorphous alloys are used as soft magnetic materials because of their high permeability and Curie temperatures. However, these alloys become brittle upon annealing, which causes serious difficulties in handling process. The annealing induced embrittlement in metallic glasses is a long standing problem, which is usually attributed to; surface or bulk crystallization [1–3], free volume annihilation [4–6], phase separation [7,8], and localized enrichment of highly diffusive elements like P, B, C etc [9]. Crystallization into brittle intermetallic compounds can straightforwardly account for the annealing embrittlement of metallic glasses [10]. However, in Fe-based amorphous alloys the embrittlement progresses well below the crystallization temperatures and the structural analysis of embrittled specimens rules out the possibility of crystallization or chemical inhomogeneity [11]. Furthermore, several nonFe-based amorphous alloys containing P, B, and/or Si do *
Corresponding author. E-mail address: kumar_golden@rediffmail.com (G. Kumar).
0022-3093/$ - see front matter 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2007.08.001
not embrittle, contradicting the viewpoint that constituents as P, B, and Si enhance the intensity of thermal embrittlement in amorphous alloys [12]. The free volume approach does not explain why the embrittlement behavior of several amorphous alloys like Fe–B and Fe–Ni–Si–B changes with a marginal change in composition where the amount of free volume is not expected to vary much [12]. Therefore, the reasons for the different thermal embrittlement behavior of Fe-based and other metallic glasses are not clear and the discrepancy arises among the proposed explanations. In crystalline materials the ductile-to-brittle transition occurs at a temperature (Tdb) above which a material is ductile and below which it breaks easily on deformation. For BCC metals, the flow stress increases with decreasing temperature because the thermal activation of dislocations is minimal, whereas the crack propagation stress is relatively independent of temperature. Thus below Tdb, the flow stress is higher than the crack propagation stress resulting in change of failure mode from ductile to brittle fracture. However, little is known about the thermal activation of shear transformation zones (STZ) which are the fundamental units of plastic deformation in metallic
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
glasses. Most of as-quenched metallic glasses show Tdb at lower than room temperature, and therefore, show good bending ductility at room temperature. Tdb in metallic glasses depends on alloy composition, quenching rate, and annealing conditions [12–14]. Annealing increases Tdb, and depending on the annealing temperature the specimen may become brittle at room temperature [12]. In the present report we correlate the annealing induced change in Tdb to the build up of atomic level stresses higher than the critical values for brittle fracture. The source of such stresses is the difference in thermal expansion of regions with different local structure. The annealing temperature after which the material becomes brittle at room temperature is referred to as the embrittling temperature Te. The Mo¨ssbauer spectra of embrittled specimens show a change in spin direction from out of plan to in-plan of the ribbons. The relation between the change in spin direction, stress state, and Te is discussed. We have chosen Fe– Ni–Si–B amorphous alloys for the present study because the Fe–Si–B amorphous alloy is a potential soft magnetic material and its thermal and magnetic properties can be tailored by a partial substitution of Fe with Ni. 2. Experimental Ingots of Fe78 xNixSi10B12 (x = 0, 19, 39, 59, 78) and Co78Si10B12 compositions were prepared by arc melting pure elements in an argon atmosphere. The composition of the as-cast ingots was verified by using chemical analysis. ribbons of 2–3 mm width and 30–40 lm thickness were prepared from the arc-melted ingots using a single roller melt spinner at a wheel speed of 40 m/s. Additionally, commercial FINEMET (Fe73.5Si13.5Nb5B7Cu1) and NANOPERM (Fe89Zr7Cu1B3) alloy ribbons obtained from Hitachi Metals were studied for comparison. The structure of the as-quenched and annealed specimens was checked by X-ray diffraction (XRD) using a Rigaku Rint 2500 (k = 0.154 nm) and transmission electron microscope (TEM) using a CM200. For the computation of pair distribution function (PDF) the electron diffraction patterns were recorded using a JEM-3010EF TEM equipped with X-type energy filter operated at 300 kV following the procedure described elsewhere [15]. The crystallization temperatures were determined from the differential scanning calorimeter (DSC) curves measured at a heating rate of 20 C/min using a Pyris-1 calorimeter. The coefficient of thermal expansion was measured with a Rigaku thermomechanical analyzer (TMA). The brittleness of the ribbons was measured by a simple bending test as discussed in previous studies [1,9]. Transmission Fe57 Mo¨ssbauer spectra were recorded at room temperature by using a conventional spectrometer equipped with a Rh 57Co source.
883
Fig. 1. The XRD patterns of melt-spun Fe78 xNixSi10B12 ribbons.
typical for an amorphous structure. Fig. 2 shows the DSC curves of as-quenched Fe78 xNixSi10B12 melt-spun ribbons. The onset temperature of crystallization (Tx) of amorphous Fe78 xNixSi10B12 ribbons decreases from 524 C for the x = 0–482 C for the x = 78 alloy. Additionally, endothermic signals corresponding to the Curie temperatures (TC) can be clearly seen at 418 C for the x = 0, 19 specimens. The as-quenched ribbons like other metallic glasses can be largely bent without forming any visible cracks. In contrary, the specimens became brittle after isochronal heating to temperatures Te, which are significantly lower than the crystallization temperatures. The Te, Tx, and TC for the different alloys are listed in Table 1. It is worth noting that Tx decreases whereas the Te increases with increasing Ni content in amorphous Fe78 xNixSi10B12 ribbons. Thus the Te and Tx values do not correlate in amorphous Fe78 xNixSi10B12 ribbons. These observations suggest that the thermal embrittlement in Fe-rich Fe78 xNixSi10B12 and Fe73.5Nb3Cu1Si13.5B9 amorphous alloys is not related to the crystallization. Fig. 3 shows the bright-field TEM
3. Results As shown in the XRD patterns (Fig. 1), the as-quenched melt-spun ribbons are amorphous displaying broad halos
Fig. 2. The DSC curves of amorphous Fe78 xNixSi10B12 melt-spun ribbons measured at a heating rate of 20 C/min.
884
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
Table 1 Crystallization temperature (Tx), embrittling temperature (Te), Curie temperature (TC), and volume magnetostriction (ks) for several Fe-based amorphous alloys Alloy
Tx (C) ± 2
Te (C) ± 2
TC (C) ± 2
(ks ± 0.1) · 10
Fe78Si10B12 Fe59Ni19Si10B12 Fe39Ni39Si10B12 Fe19Ni59Si10B12 Ni78Si10B12 Co78Si10B12 Fe73.5Nb3Cu1Si13.5B9 Fe89Zr7Cu1B3
524 500 482 510 440 420 515 500
300 310 335 362 455 385 300 475
418 418
27
313 <50
6
2.1 1.1
images and the corresponding selected area electron diffraction (SAED) patterns for the embrittled Fe59Ni19Si10B12 (Fig. 3(a)) and Fe39Ni39Si10B12 (Fig. 3(b)) ribbons. The featureless TEM images and broad halos in the SAED patterns confirm the amorphous structure of the embrittled ribbons. Similar results were obtained for the other embrittled specimens except Ni78Si10B12, which showed the formation of FCC-Ni crystals after heating to 450 C. Pair distribution function g(r) profiles were deduced from the intensity of zero-loss SAED patterns in order to analyze the effect of annealing on short-range order (SRO). The details of g(r) calculation are given elsewhere [15]. Fig. 4 shows the g(r) profiles of Fe59Ni19Si10B12 melt-spun ribbons in as-quenched state and after isochronal annealing. The shape of main peak in the g(r) of asquenched sample is symmetric and there are no features that can be termed as pre-peak or splitting of the first peak. There is no discernible change in the main g(r) peak of the specimens annealed at 300 C (still ductile) and 400 C (brittle) except small increase in intensity. This indicates that there is no detectable change in the short-range order associated with the thermal embrittlement in amorphous Fe59Ni19Si10B12 ribbons. Similar results were obtained for the other Fe78 xNixSi10B12 (0 6 x 6 59) amorphous ribbons.
Fig. 4. Pair distribution functions g(r) of as-quenched and annealed Fe59Ni19Si10B12 melt-spun ribbons.
Fig. 5(a) shows the fractional change in length as a function of temperature for the amorphous Fe78 xNixSi10B12, Co78Si10B12, Fe89Zr7Cu1B3, and Fe73.5Nb3Cu1Si13.5B9 ribbons. The Fe78 xNixSi10B12 (x = 0, 19, 39) and Fe73.5Nb3Cu1Si13.5B9 ribbons show a linear increase in length in the beginning followed by a flat region (lower thermal expansion) close to their Curie temperatures. This flat region in the TMA curves is related to the Invar characteristic caused by large positive spontaneous volume magnetostriction of Fe-based amorphous alloys [16]. A notable feature in the TMA curves is the absence of Invar-type non-linearity in the thermal expansion of amorphous Ni78Si10B12, Co78Si10B12, and Fe89Zr7Cu1B3 ribbons. Interestingly, these alloys do not become brittle after annealing below the crystallization temperatures. As marked by arrows, the embrittling temperature for most of the amorphous Fe78 xNixSi10B12 ribbons falls in the range of Invar region. Fig. 5(b) shows the variation of thermal expansion coefficients calculated from the TMA curves shown in Fig. 5(a). The value for BCC-Fe is also shown for comparison as the local SRO in Fe-based amorphous alloys is
Fig. 3. Bright-field TEM images and the corresponding SAED patterns of embrittled Fe59Ni19Si10B12 (a) Fe39Ni39Si10B12 and (b) melt-spun ribbons.
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
885
about the change in magnetic environment of Fe atoms. Fig. 6(a) shows the Mo¨ssbauer spectra of amorphous Fe59Ni19Si10B12 ribbon in the as-quenched (ductile) state and after annealing 400 C (brittle). In each case, the spectrum is composed of six broadened lines, which is characteristic of amorphous ferromagnetic alloys. The spectra were fitted to six Lorentzian peaks, characterized by a mean hyperfine field Heff and a mean quadrupole interaction. The fitted results, shown in Fig. 6(a) as the solid lines, describe the data satisfactorily. Fig. 6(b) shows the distribution P(H) of hyperfine field calculated from the fitting of the Mo¨ssbauer spectra. The P(H) shows a symmetric peak at about 246 kOe for the as-quenched as well as the annealed specimen. The Mo¨ssbauer analysis indicates that there is no long-range structural change associated with the annealing up to 400 C. Though, the spectra appear similar before and after annealing, there is a change in the relative intensity or area of the spectral lines. For a thin sample the area ratios of the six lines are A1,6:A2,5:A3,4::3:A2,5:1. Here, A2,5 = 4sin2 //(1 + cos2 /) and / is the angle between the spin orientation and the c-ray direction. Thus using the ratio A2,5/A3,4 in the Mo¨ssbauer spectra, the variation in
Fig. 5. (a) The fractional change in length and (b) thermal expansion coefficient of Fe-based amorphous melt-spun ribbons compared with a Co-based amorphous alloy and BCC-Fe. The arrows indicate the embrittling temperatures.
expected to be like BCC-Fe. The thermal expansion of amorphous Fe78 xNixSi10B12 (x < 78) ribbons is lower than that of BCC-Fe. The difference in thermal expansion between BCC-Fe and amorphous Fe78 xNixSi10B12 ribbons decreases with increasing Ni content. The thermal expansion coefficient of Fe78 xNixSi10B12 ribbons shows a significant (20–40%) reduction close to the Curie temperatures. The observed reduction in the thermal expansion of amorphous Fe78 xNixSi10B12 alloys as a function of temperature (under a small tensile stress) is related to their Invar characteristics. This change is not caused by the reduction of free volume because the extent of thermal expansion reduction did not seem to change with the quenching rate. Furthermore, the Co78Si10B12 and Fe89Zr7Cu1B3 ribbons do not show decrease in thermal expansion due to the absence of Invar effect. A possible relation between the Invar effect and the thermal embrittlement of Fe-based amorphous alloys will be discussed in the next section. The Mo¨ssbauer spectra were recorded as a function of annealing temperature in order to gain information
Fig. 6. (a) Mo¨ssbauer spectra of Fe59Ni19Si10B12 melt-spun ribbon in the as-quenched state and after annealing at 400 C. (b) The distribution of hyperfine field calculated from the fitted Mo¨ssbauer spectra.
886
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
Fig. 7. The area ratio (A2,5/A3,4) as a function of annealing temperature derived from the Mo¨ssbauer spectra of Fe59Ni19Si10B12 melt-spun ribbon.
spin orientations with increasing annealing temperature can be obtained. The ratio A2,5/A3,4 increases from 0 to 4 as / changes from 0 (all moments parallel to the c-ray direction) to 90 (all moments perpendicular to the c-ray direction). For a completely random spin orientations the A2,5/A3,4 ratio is 2. Fig. 7 shows the A2,5/A3,4 ratio as a function of annealing temperature. The incident c-ray was perpendicular to the ribbon plan. In the analysis, it was assumed that the ratio A1,6/A3,4 is 3 and the ratio A2,5/A3,4 was determined with a common fitting parameter for the area of six components. The calculated A2,5/A3,4 ratio remains unchanged for the as-quenched and the ribbons annealed up to 275 C. However, it increases by about 20% for the specimens annealed at 400 C and 450 C. This indicates that at temperatures higher than 275 C there is change in the spin orientations into the ribbon plane. The A2,5/A3,4 ratio for the as-quenched ribbon is about 2.9, which suggests that the spin orientations are neither completely random nor in-plan of the ribbon. The directions of magnetic moments in the as-prepared ferromagnetic amorphous alloys are determined by the quenched-in strain fields. It is therefore not surprising that some temperature dependence of magnetization should occur and change in spin directions as a function of annealing has been previously observed for several Fe-based amorphous alloys [17,18]. However, the notable observation in the present results is that the change in spin orientations with annealing temperature shows a sudden transition that matches closely with the embrittling temperature. These results indicate a possible relationship between the change in stress state and the embrittling temperature as discussed in the following section. 4. Discussion Despite containing the same metalloid concentrations, the thermal embrittlement behavior of amorphous Fe78Si10B12, Ni78Si10B12, and Co78Si10B12 ribbons is signif-
icantly different. This suggests that metalloid enrichment may not be the dominating factor for thermal embrittlement in the above mentioned alloys. Inoue et al. reported the change in crystallization product from BCC-Fe for Fe78Si10B12 to FCC-Ni for Ni78Si10B12 amorphous alloys as the cause for their different embrittlement behavior [1]. However, amorphous Fe89Zr7Cu1B3 alloy with BCC-Fe crystallization product does not become brittle at room temperature after annealing below the crystallization temperature. Furthermore, the possibility of crystallization can be excluded on the basis of the X-ray and TEM analysis, which show that the embrittled specimens are still amorphous. Recent TEM results have suggested the existence of BCC-type local SRO in Fe-based amorphous alloys [19]. The Invar effect and thermal embrittlement in Fe-based amorphous alloys can be explained by assuming the presence of BCC-type ordered regions embedded in the amorphous matrix. The matrix with Invar characteristics shows a small thermal expansion whereas the ordered regions exhibit thermal expansion similar to BCC-Fe. This difference in thermal expansion can result in thermal stresses in the ribbons during cooling after annealing. The ribbons will become brittle when the thermal stress becomes higher than the critical fracture stress. The magnitude of thermal stress will depend on the annealing temperature and the difference in the thermal expansion of ordered regions and matrix. As shown in Fig. 5(a), the increases in Te with increasing Ni content of amorphous Fe78 xNixSi10B12 ribbons can be explained by the smaller difference in thermal expansion of the ribbon and BCC-Fe. The Febased amorphous alloys without Invar effect are structurally more homogeneous and, therefore, more resistant to the thermal embrittlement. This hypothesis of different regions with different thermal expansions can also explain the previous results where it was shown that the amorphous alloys with tendency to phase separate show pronounced thermal embrittlement. Therefore, the key factor for the thermal embrittlement of amorphous ribbons appears to be the stress caused by structural, compositional or any kind of inhomogeneity. Mo¨ssbauer spectra of the as-quenched (ductile) and annealed (brittle) ribbons show that the transition from ductile to brittle fracture is accompanied by a relative change in the magnetization directions from out-of-plan to in-plan of the ribbon. The strain pattern in the asquenched ribbons is rather complex, which depends on material and preparation conditions. Therefore, the spin directions in the as-quenched ribbons may vary depending on the quenching conditions. However, the quenched-in strains are relaxed upon annealing, and during subsequent cooling the competition between magnetic anisotropy and residual stresses will determine the spin orientations. The Mo¨ssbauer analysis reveals that majority of the spins are oriented in-plan after heat treatment above certain critical temperature. One source of this change in spin directions is the shape anisotropy, which dominates after the relaxation
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
887
a
b Matrix after accommodation BCC-type Cluster
ΔL/L
Matrix Source of residual stress
Temperature Fig. 8. A schematic illustration of amorphous structure for Fe-based ribbons displaying BCC-type clusters embedded in matrix with positive magnetostriction.
of quenched-in strains. Another reason for change in spin direction after annealing may be the build up of tensile stress along the ribbon plan during cooling. The magnetization direction will be favorable along the tensile stress direction as the majority of the alloys studied here show positive magnetostriction. This tensile stress build up as a result of different thermal expansion regions can be a crucial factor for the brittleness of annealed amorphous alloys. The overall scenario can be explained by a schematic diagram illustrated in Fig. 8. The structure of Fe-based amorphous alloys can be described as a mixture of BCCtype ordered clusters randomly distributed in matrix (Fig. 8(a)). The matrix exhibits positive magnetostriction whereas the BCC-type clusters behave like crystalline Fe. Therefore, the length change of matrix follows broken curve whereas the BCC-type clusters follow dotted line in Fig. 8(b). During annealing, the volume expansion of BCC-type clusters is accommodated by viscous flow of the matrix. However, the subsequent cooling would result in a residual tensile stress at the BCC-type clusters due to the low thermal expansion of the matrix, which is schematically shown as solid curve in Fig. 8(b). The amount of ten-
sile stress retained at room temperature depends on the annealing temperature. With increasing Ni content in amorphous Fe78 xNixSi10B12 ribbons, the Invar effect lessens or in other words its thermal expansion increases. This reduces the thermal stress and consequently raises the embrittling temperature. 5. Conclusion The structural analysis revealed no detectable change in structure or composition associated with the early thermal embrittling of several Fe-based amorphous alloys. However, the thermal expansion measurements and the Mo¨ssbauer results gave valuable insights about the possible reasons for the thermal embrittlement. The Fe-based amorphous alloys showing Invar effect with a small thermal expansion are more prone to the thermal embrittlement. The Mo¨ssbauer analysis reveals that the magnetization direction changes from out-of-plan to in-plan of the annealed (brittle) ribbons. This transition of magnetization direction appears to be related to the change in stress state. A build up of tensile stress can be envisaged from the
888
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
positive (volume) magnetostriction values. While it is difficult to draw any quantitative conclusion from such a comparison, it may prove instructive to recall a change in stress state because of annealing and subsequent cooling as the possible reason for the thermal embrittlement of Febased glasses. The factors which contribute to the stress development are the fluctuations of thermal expansion stemming from structural, compositional, and magnetic inhomogeneities. Acknowledgment This work was supported by the Grant in Aid of Ministry of Education, Sports, Culture, Science and Technology, Priority Area on ‘Materials Science of Bulk Metallic Glasses’. References [1] A. Inoue, T. Masumoto, H. Kimura, Sci. Rep. RITU 27 (1979) 159.
[2] P. Lamparter, S. Steeb, J. Non-Cryst. Solids 106 (1988) 137. [3] C. Janot, B. George, D. Teirlinck, G. Marchal, C. Tete, P. Delcroix, Philos. Mag. 47 (1983) 301. [4] D. Deng, A.S. Argon, Acta Mater. 34 (1986) 2011. [5] R. Gerling, F.P. Schimansky, R. Wagner, Acta Mater. 36 (1988) 575. [6] T.W. Wu, F. Spaepen, Philos. Mag. 61 (1990) 739. [7] H.S. Chen, Scripta Mater. 11 (1977) 367. [8] J. Piller, P. Haasen, Acta Mater. 30 (1982) 1. [9] F.E. Luborsky, J.L. Walter, J. Appl. Phys. 47 (1976) 3648. [10] Y.C. Niu, X.F. Bian, W.M. Wang, J. Non-Cryst. Solids 341 (2004) 40. [11] G.E. Fish, H.R. Child, J. Appl. Phys. 52 (1981) 1880. [12] A.R. Yavari, Mater. Sci. Eng. 98 (1988) 491. [13] G.C. Chi, H.S. Chen, C.E. Miller, J. Appl. Phys. 49 (1978) 1715. [14] P. Allia, P. Tiberto, M. Baricco, F. Vinai, Appl. Phys. Lett. 63 (1993) 2759. [15] T. Ohkubo, Y. Hirotsu, Phys. Rev. B 67 (2003) 094201. [16] M. Takahashi, M. Koshimura, T. Abuzuka, Jpn. J. Appl. Phys. 20 (1981) 1821. [17] C.L. Chien, Phys. Rev. B 18 (1978) 1003. [18] H.N. Ok, A.H. Morrish, Phys. Rev. B 22 (1980) 4215. [19] T. Ohkubo, Y. Hirotsu, Mater. Sci. Eng. A 304–306 (2001) 300.
Journal of Non-Crystalline Solids 354 (2008) 882–888 www.elsevier.com/locate/jnoncrysol
Thermal embrittlement of Fe-based amorphous ribbons G. Kumar *, M. Ohnuma, T. Furubayashi, T. Ohkubo, K. Hono National Institute for Materials Science, 1-2-1 Sengen, Tsukuba 305-0047, Japan Received 6 April 2007; received in revised form 31 July 2007 Available online 14 September 2007
Abstract Thermal embrittlement of amorphous Fe78 xNixSi10B12 and Co78Si10B12 alloys has been studied. The majority of Fe-based amorphous alloys became brittle after annealing to temperatures significantly lower than their crystallization temperatures, whereas the non-Fe-based amorphous alloys remained ductile after annealing up to their crystallization temperatures. The present results emphasize on the key difference between the embrittlement behavior of Fe-based and non-Fe-based amorphous alloys. We point out a possible correlation between different phenomena like Invar effect, compositional inhomogenity, and residual stress playing a critical role in the thermal embrittlement of Fe-based amorphous alloys. 2007 Elsevier B.V. All rights reserved. PACS: 65.60.+a; 75.50.Bb; 81.05.Kf Keywords: Amorphous metals, metallic glasses; Transition metals; Mechanical, stress relaxation; Short-range order
1. Introduction Fe-based amorphous alloys are used as soft magnetic materials because of their high permeability and Curie temperatures. However, these alloys become brittle upon annealing, which causes serious difficulties in handling process. The annealing induced embrittlement in metallic glasses is a long standing problem, which is usually attributed to; surface or bulk crystallization [1–3], free volume annihilation [4–6], phase separation [7,8], and localized enrichment of highly diffusive elements like P, B, C etc [9]. Crystallization into brittle intermetallic compounds can straightforwardly account for the annealing embrittlement of metallic glasses [10]. However, in Fe-based amorphous alloys the embrittlement progresses well below the crystallization temperatures and the structural analysis of embrittled specimens rules out the possibility of crystallization or chemical inhomogeneity [11]. Furthermore, several nonFe-based amorphous alloys containing P, B, and/or Si do *
Corresponding author. E-mail address: kumar_golden@rediffmail.com (G. Kumar).
0022-3093/$ - see front matter 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2007.08.001
not embrittle, contradicting the viewpoint that constituents as P, B, and Si enhance the intensity of thermal embrittlement in amorphous alloys [12]. The free volume approach does not explain why the embrittlement behavior of several amorphous alloys like Fe–B and Fe–Ni–Si–B changes with a marginal change in composition where the amount of free volume is not expected to vary much [12]. Therefore, the reasons for the different thermal embrittlement behavior of Fe-based and other metallic glasses are not clear and the discrepancy arises among the proposed explanations. In crystalline materials the ductile-to-brittle transition occurs at a temperature (Tdb) above which a material is ductile and below which it breaks easily on deformation. For BCC metals, the flow stress increases with decreasing temperature because the thermal activation of dislocations is minimal, whereas the crack propagation stress is relatively independent of temperature. Thus below Tdb, the flow stress is higher than the crack propagation stress resulting in change of failure mode from ductile to brittle fracture. However, little is known about the thermal activation of shear transformation zones (STZ) which are the fundamental units of plastic deformation in metallic
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
glasses. Most of as-quenched metallic glasses show Tdb at lower than room temperature, and therefore, show good bending ductility at room temperature. Tdb in metallic glasses depends on alloy composition, quenching rate, and annealing conditions [12–14]. Annealing increases Tdb, and depending on the annealing temperature the specimen may become brittle at room temperature [12]. In the present report we correlate the annealing induced change in Tdb to the build up of atomic level stresses higher than the critical values for brittle fracture. The source of such stresses is the difference in thermal expansion of regions with different local structure. The annealing temperature after which the material becomes brittle at room temperature is referred to as the embrittling temperature Te. The Mo¨ssbauer spectra of embrittled specimens show a change in spin direction from out of plan to in-plan of the ribbons. The relation between the change in spin direction, stress state, and Te is discussed. We have chosen Fe– Ni–Si–B amorphous alloys for the present study because the Fe–Si–B amorphous alloy is a potential soft magnetic material and its thermal and magnetic properties can be tailored by a partial substitution of Fe with Ni. 2. Experimental Ingots of Fe78 xNixSi10B12 (x = 0, 19, 39, 59, 78) and Co78Si10B12 compositions were prepared by arc melting pure elements in an argon atmosphere. The composition of the as-cast ingots was verified by using chemical analysis. ribbons of 2–3 mm width and 30–40 lm thickness were prepared from the arc-melted ingots using a single roller melt spinner at a wheel speed of 40 m/s. Additionally, commercial FINEMET (Fe73.5Si13.5Nb5B7Cu1) and NANOPERM (Fe89Zr7Cu1B3) alloy ribbons obtained from Hitachi Metals were studied for comparison. The structure of the as-quenched and annealed specimens was checked by X-ray diffraction (XRD) using a Rigaku Rint 2500 (k = 0.154 nm) and transmission electron microscope (TEM) using a CM200. For the computation of pair distribution function (PDF) the electron diffraction patterns were recorded using a JEM-3010EF TEM equipped with X-type energy filter operated at 300 kV following the procedure described elsewhere [15]. The crystallization temperatures were determined from the differential scanning calorimeter (DSC) curves measured at a heating rate of 20 C/min using a Pyris-1 calorimeter. The coefficient of thermal expansion was measured with a Rigaku thermomechanical analyzer (TMA). The brittleness of the ribbons was measured by a simple bending test as discussed in previous studies [1,9]. Transmission Fe57 Mo¨ssbauer spectra were recorded at room temperature by using a conventional spectrometer equipped with a Rh 57Co source.
883
Fig. 1. The XRD patterns of melt-spun Fe78 xNixSi10B12 ribbons.
typical for an amorphous structure. Fig. 2 shows the DSC curves of as-quenched Fe78 xNixSi10B12 melt-spun ribbons. The onset temperature of crystallization (Tx) of amorphous Fe78 xNixSi10B12 ribbons decreases from 524 C for the x = 0–482 C for the x = 78 alloy. Additionally, endothermic signals corresponding to the Curie temperatures (TC) can be clearly seen at 418 C for the x = 0, 19 specimens. The as-quenched ribbons like other metallic glasses can be largely bent without forming any visible cracks. In contrary, the specimens became brittle after isochronal heating to temperatures Te, which are significantly lower than the crystallization temperatures. The Te, Tx, and TC for the different alloys are listed in Table 1. It is worth noting that Tx decreases whereas the Te increases with increasing Ni content in amorphous Fe78 xNixSi10B12 ribbons. Thus the Te and Tx values do not correlate in amorphous Fe78 xNixSi10B12 ribbons. These observations suggest that the thermal embrittlement in Fe-rich Fe78 xNixSi10B12 and Fe73.5Nb3Cu1Si13.5B9 amorphous alloys is not related to the crystallization. Fig. 3 shows the bright-field TEM
3. Results As shown in the XRD patterns (Fig. 1), the as-quenched melt-spun ribbons are amorphous displaying broad halos
Fig. 2. The DSC curves of amorphous Fe78 xNixSi10B12 melt-spun ribbons measured at a heating rate of 20 C/min.
884
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
Table 1 Crystallization temperature (Tx), embrittling temperature (Te), Curie temperature (TC), and volume magnetostriction (ks) for several Fe-based amorphous alloys Alloy
Tx (C) ± 2
Te (C) ± 2
TC (C) ± 2
(ks ± 0.1) · 10
Fe78Si10B12 Fe59Ni19Si10B12 Fe39Ni39Si10B12 Fe19Ni59Si10B12 Ni78Si10B12 Co78Si10B12 Fe73.5Nb3Cu1Si13.5B9 Fe89Zr7Cu1B3
524 500 482 510 440 420 515 500
300 310 335 362 455 385 300 475
418 418
27
313 <50
6
2.1 1.1
images and the corresponding selected area electron diffraction (SAED) patterns for the embrittled Fe59Ni19Si10B12 (Fig. 3(a)) and Fe39Ni39Si10B12 (Fig. 3(b)) ribbons. The featureless TEM images and broad halos in the SAED patterns confirm the amorphous structure of the embrittled ribbons. Similar results were obtained for the other embrittled specimens except Ni78Si10B12, which showed the formation of FCC-Ni crystals after heating to 450 C. Pair distribution function g(r) profiles were deduced from the intensity of zero-loss SAED patterns in order to analyze the effect of annealing on short-range order (SRO). The details of g(r) calculation are given elsewhere [15]. Fig. 4 shows the g(r) profiles of Fe59Ni19Si10B12 melt-spun ribbons in as-quenched state and after isochronal annealing. The shape of main peak in the g(r) of asquenched sample is symmetric and there are no features that can be termed as pre-peak or splitting of the first peak. There is no discernible change in the main g(r) peak of the specimens annealed at 300 C (still ductile) and 400 C (brittle) except small increase in intensity. This indicates that there is no detectable change in the short-range order associated with the thermal embrittlement in amorphous Fe59Ni19Si10B12 ribbons. Similar results were obtained for the other Fe78 xNixSi10B12 (0 6 x 6 59) amorphous ribbons.
Fig. 4. Pair distribution functions g(r) of as-quenched and annealed Fe59Ni19Si10B12 melt-spun ribbons.
Fig. 5(a) shows the fractional change in length as a function of temperature for the amorphous Fe78 xNixSi10B12, Co78Si10B12, Fe89Zr7Cu1B3, and Fe73.5Nb3Cu1Si13.5B9 ribbons. The Fe78 xNixSi10B12 (x = 0, 19, 39) and Fe73.5Nb3Cu1Si13.5B9 ribbons show a linear increase in length in the beginning followed by a flat region (lower thermal expansion) close to their Curie temperatures. This flat region in the TMA curves is related to the Invar characteristic caused by large positive spontaneous volume magnetostriction of Fe-based amorphous alloys [16]. A notable feature in the TMA curves is the absence of Invar-type non-linearity in the thermal expansion of amorphous Ni78Si10B12, Co78Si10B12, and Fe89Zr7Cu1B3 ribbons. Interestingly, these alloys do not become brittle after annealing below the crystallization temperatures. As marked by arrows, the embrittling temperature for most of the amorphous Fe78 xNixSi10B12 ribbons falls in the range of Invar region. Fig. 5(b) shows the variation of thermal expansion coefficients calculated from the TMA curves shown in Fig. 5(a). The value for BCC-Fe is also shown for comparison as the local SRO in Fe-based amorphous alloys is
Fig. 3. Bright-field TEM images and the corresponding SAED patterns of embrittled Fe59Ni19Si10B12 (a) Fe39Ni39Si10B12 and (b) melt-spun ribbons.
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
885
about the change in magnetic environment of Fe atoms. Fig. 6(a) shows the Mo¨ssbauer spectra of amorphous Fe59Ni19Si10B12 ribbon in the as-quenched (ductile) state and after annealing 400 C (brittle). In each case, the spectrum is composed of six broadened lines, which is characteristic of amorphous ferromagnetic alloys. The spectra were fitted to six Lorentzian peaks, characterized by a mean hyperfine field Heff and a mean quadrupole interaction. The fitted results, shown in Fig. 6(a) as the solid lines, describe the data satisfactorily. Fig. 6(b) shows the distribution P(H) of hyperfine field calculated from the fitting of the Mo¨ssbauer spectra. The P(H) shows a symmetric peak at about 246 kOe for the as-quenched as well as the annealed specimen. The Mo¨ssbauer analysis indicates that there is no long-range structural change associated with the annealing up to 400 C. Though, the spectra appear similar before and after annealing, there is a change in the relative intensity or area of the spectral lines. For a thin sample the area ratios of the six lines are A1,6:A2,5:A3,4::3:A2,5:1. Here, A2,5 = 4sin2 //(1 + cos2 /) and / is the angle between the spin orientation and the c-ray direction. Thus using the ratio A2,5/A3,4 in the Mo¨ssbauer spectra, the variation in
Fig. 5. (a) The fractional change in length and (b) thermal expansion coefficient of Fe-based amorphous melt-spun ribbons compared with a Co-based amorphous alloy and BCC-Fe. The arrows indicate the embrittling temperatures.
expected to be like BCC-Fe. The thermal expansion of amorphous Fe78 xNixSi10B12 (x < 78) ribbons is lower than that of BCC-Fe. The difference in thermal expansion between BCC-Fe and amorphous Fe78 xNixSi10B12 ribbons decreases with increasing Ni content. The thermal expansion coefficient of Fe78 xNixSi10B12 ribbons shows a significant (20–40%) reduction close to the Curie temperatures. The observed reduction in the thermal expansion of amorphous Fe78 xNixSi10B12 alloys as a function of temperature (under a small tensile stress) is related to their Invar characteristics. This change is not caused by the reduction of free volume because the extent of thermal expansion reduction did not seem to change with the quenching rate. Furthermore, the Co78Si10B12 and Fe89Zr7Cu1B3 ribbons do not show decrease in thermal expansion due to the absence of Invar effect. A possible relation between the Invar effect and the thermal embrittlement of Fe-based amorphous alloys will be discussed in the next section. The Mo¨ssbauer spectra were recorded as a function of annealing temperature in order to gain information
Fig. 6. (a) Mo¨ssbauer spectra of Fe59Ni19Si10B12 melt-spun ribbon in the as-quenched state and after annealing at 400 C. (b) The distribution of hyperfine field calculated from the fitted Mo¨ssbauer spectra.
886
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
Fig. 7. The area ratio (A2,5/A3,4) as a function of annealing temperature derived from the Mo¨ssbauer spectra of Fe59Ni19Si10B12 melt-spun ribbon.
spin orientations with increasing annealing temperature can be obtained. The ratio A2,5/A3,4 increases from 0 to 4 as / changes from 0 (all moments parallel to the c-ray direction) to 90 (all moments perpendicular to the c-ray direction). For a completely random spin orientations the A2,5/A3,4 ratio is 2. Fig. 7 shows the A2,5/A3,4 ratio as a function of annealing temperature. The incident c-ray was perpendicular to the ribbon plan. In the analysis, it was assumed that the ratio A1,6/A3,4 is 3 and the ratio A2,5/A3,4 was determined with a common fitting parameter for the area of six components. The calculated A2,5/A3,4 ratio remains unchanged for the as-quenched and the ribbons annealed up to 275 C. However, it increases by about 20% for the specimens annealed at 400 C and 450 C. This indicates that at temperatures higher than 275 C there is change in the spin orientations into the ribbon plane. The A2,5/A3,4 ratio for the as-quenched ribbon is about 2.9, which suggests that the spin orientations are neither completely random nor in-plan of the ribbon. The directions of magnetic moments in the as-prepared ferromagnetic amorphous alloys are determined by the quenched-in strain fields. It is therefore not surprising that some temperature dependence of magnetization should occur and change in spin directions as a function of annealing has been previously observed for several Fe-based amorphous alloys [17,18]. However, the notable observation in the present results is that the change in spin orientations with annealing temperature shows a sudden transition that matches closely with the embrittling temperature. These results indicate a possible relationship between the change in stress state and the embrittling temperature as discussed in the following section. 4. Discussion Despite containing the same metalloid concentrations, the thermal embrittlement behavior of amorphous Fe78Si10B12, Ni78Si10B12, and Co78Si10B12 ribbons is signif-
icantly different. This suggests that metalloid enrichment may not be the dominating factor for thermal embrittlement in the above mentioned alloys. Inoue et al. reported the change in crystallization product from BCC-Fe for Fe78Si10B12 to FCC-Ni for Ni78Si10B12 amorphous alloys as the cause for their different embrittlement behavior [1]. However, amorphous Fe89Zr7Cu1B3 alloy with BCC-Fe crystallization product does not become brittle at room temperature after annealing below the crystallization temperature. Furthermore, the possibility of crystallization can be excluded on the basis of the X-ray and TEM analysis, which show that the embrittled specimens are still amorphous. Recent TEM results have suggested the existence of BCC-type local SRO in Fe-based amorphous alloys [19]. The Invar effect and thermal embrittlement in Fe-based amorphous alloys can be explained by assuming the presence of BCC-type ordered regions embedded in the amorphous matrix. The matrix with Invar characteristics shows a small thermal expansion whereas the ordered regions exhibit thermal expansion similar to BCC-Fe. This difference in thermal expansion can result in thermal stresses in the ribbons during cooling after annealing. The ribbons will become brittle when the thermal stress becomes higher than the critical fracture stress. The magnitude of thermal stress will depend on the annealing temperature and the difference in the thermal expansion of ordered regions and matrix. As shown in Fig. 5(a), the increases in Te with increasing Ni content of amorphous Fe78 xNixSi10B12 ribbons can be explained by the smaller difference in thermal expansion of the ribbon and BCC-Fe. The Febased amorphous alloys without Invar effect are structurally more homogeneous and, therefore, more resistant to the thermal embrittlement. This hypothesis of different regions with different thermal expansions can also explain the previous results where it was shown that the amorphous alloys with tendency to phase separate show pronounced thermal embrittlement. Therefore, the key factor for the thermal embrittlement of amorphous ribbons appears to be the stress caused by structural, compositional or any kind of inhomogeneity. Mo¨ssbauer spectra of the as-quenched (ductile) and annealed (brittle) ribbons show that the transition from ductile to brittle fracture is accompanied by a relative change in the magnetization directions from out-of-plan to in-plan of the ribbon. The strain pattern in the asquenched ribbons is rather complex, which depends on material and preparation conditions. Therefore, the spin directions in the as-quenched ribbons may vary depending on the quenching conditions. However, the quenched-in strains are relaxed upon annealing, and during subsequent cooling the competition between magnetic anisotropy and residual stresses will determine the spin orientations. The Mo¨ssbauer analysis reveals that majority of the spins are oriented in-plan after heat treatment above certain critical temperature. One source of this change in spin directions is the shape anisotropy, which dominates after the relaxation
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
887
a
b Matrix after accommodation BCC-type Cluster
ΔL/L
Matrix Source of residual stress
Temperature Fig. 8. A schematic illustration of amorphous structure for Fe-based ribbons displaying BCC-type clusters embedded in matrix with positive magnetostriction.
of quenched-in strains. Another reason for change in spin direction after annealing may be the build up of tensile stress along the ribbon plan during cooling. The magnetization direction will be favorable along the tensile stress direction as the majority of the alloys studied here show positive magnetostriction. This tensile stress build up as a result of different thermal expansion regions can be a crucial factor for the brittleness of annealed amorphous alloys. The overall scenario can be explained by a schematic diagram illustrated in Fig. 8. The structure of Fe-based amorphous alloys can be described as a mixture of BCCtype ordered clusters randomly distributed in matrix (Fig. 8(a)). The matrix exhibits positive magnetostriction whereas the BCC-type clusters behave like crystalline Fe. Therefore, the length change of matrix follows broken curve whereas the BCC-type clusters follow dotted line in Fig. 8(b). During annealing, the volume expansion of BCC-type clusters is accommodated by viscous flow of the matrix. However, the subsequent cooling would result in a residual tensile stress at the BCC-type clusters due to the low thermal expansion of the matrix, which is schematically shown as solid curve in Fig. 8(b). The amount of ten-
sile stress retained at room temperature depends on the annealing temperature. With increasing Ni content in amorphous Fe78 xNixSi10B12 ribbons, the Invar effect lessens or in other words its thermal expansion increases. This reduces the thermal stress and consequently raises the embrittling temperature. 5. Conclusion The structural analysis revealed no detectable change in structure or composition associated with the early thermal embrittling of several Fe-based amorphous alloys. However, the thermal expansion measurements and the Mo¨ssbauer results gave valuable insights about the possible reasons for the thermal embrittlement. The Fe-based amorphous alloys showing Invar effect with a small thermal expansion are more prone to the thermal embrittlement. The Mo¨ssbauer analysis reveals that the magnetization direction changes from out-of-plan to in-plan of the annealed (brittle) ribbons. This transition of magnetization direction appears to be related to the change in stress state. A build up of tensile stress can be envisaged from the
888
G. Kumar et al. / Journal of Non-Crystalline Solids 354 (2008) 882–888
positive (volume) magnetostriction values. While it is difficult to draw any quantitative conclusion from such a comparison, it may prove instructive to recall a change in stress state because of annealing and subsequent cooling as the possible reason for the thermal embrittlement of Febased glasses. The factors which contribute to the stress development are the fluctuations of thermal expansion stemming from structural, compositional, and magnetic inhomogeneities. Acknowledgment This work was supported by the Grant in Aid of Ministry of Education, Sports, Culture, Science and Technology, Priority Area on ‘Materials Science of Bulk Metallic Glasses’. References [1] A. Inoue, T. Masumoto, H. Kimura, Sci. Rep. RITU 27 (1979) 159.
[2] P. Lamparter, S. Steeb, J. Non-Cryst. Solids 106 (1988) 137. [3] C. Janot, B. George, D. Teirlinck, G. Marchal, C. Tete, P. Delcroix, Philos. Mag. 47 (1983) 301. [4] D. Deng, A.S. Argon, Acta Mater. 34 (1986) 2011. [5] R. Gerling, F.P. Schimansky, R. Wagner, Acta Mater. 36 (1988) 575. [6] T.W. Wu, F. Spaepen, Philos. Mag. 61 (1990) 739. [7] H.S. Chen, Scripta Mater. 11 (1977) 367. [8] J. Piller, P. Haasen, Acta Mater. 30 (1982) 1. [9] F.E. Luborsky, J.L. Walter, J. Appl. Phys. 47 (1976) 3648. [10] Y.C. Niu, X.F. Bian, W.M. Wang, J. Non-Cryst. Solids 341 (2004) 40. [11] G.E. Fish, H.R. Child, J. Appl. Phys. 52 (1981) 1880. [12] A.R. Yavari, Mater. Sci. Eng. 98 (1988) 491. [13] G.C. Chi, H.S. Chen, C.E. Miller, J. Appl. Phys. 49 (1978) 1715. [14] P. Allia, P. Tiberto, M. Baricco, F. Vinai, Appl. Phys. Lett. 63 (1993) 2759. [15] T. Ohkubo, Y. Hirotsu, Phys. Rev. B 67 (2003) 094201. [16] M. Takahashi, M. Koshimura, T. Abuzuka, Jpn. J. Appl. Phys. 20 (1981) 1821. [17] C.L. Chien, Phys. Rev. B 18 (1978) 1003. [18] H.N. Ok, A.H. Morrish, Phys. Rev. B 22 (1980) 4215. [19] T. Ohkubo, Y. Hirotsu, Mater. Sci. Eng. A 304–306 (2001) 300.