Glass-forming ability of melt-spun multicomponent (Ti, Zr, Hf) (Cu, Ni, Co) Al alloys with equiatomic substitution

Transcription

1 Journal of Non-Crystalline Solids 347 (2004) Glass-forming ability of melt-spun multicomponent (Ti, Zr, Hf) (Cu, Ni, Co) Al alloys with equiatomic substitution L.C. Zhang, J. Xu * Shenyang National Laboratory for Materials Science, Institute of Metal Research, CAS, 72 Wenhua Road, Shenyang , PeopleÕs Republic of China Received 22 February 2004; received in revised form 3 June 2004 Abstract (Ti 0.33 Zr 0.33 Hf 0.33 ) x (Ni 0.33 Cu 0.33 Co 0.33 ) 90 x Al 10 (x = 50, 55, 60) multicomponent alloys with equiatomic substitution for early and late transition metal contents were rapidly quenched by melt-spinning with the wheel speeds of 5 39m/s. Structural features of the as-quenched ribbons were characterized using X-ray diffraction, transmission electron microscopy and differential scanning calorimetry. Using the wheel speeds higher than 30m/s, fully glassy ribbons are obtained for all three alloys. With increasing content of the early transition metals, the glass-forming ability of the alloys was degraded. For the best glass-forming alloy at x = 50, the maximum thickness of fully glassy alloy ribbons reached about 85lm. The width of supercooled liquid region, DT x and the reduced glass transition temperature, T rg for this glassy alloy was determined to be 124K and 0.47, respectively. Compared with the (Ti, Zr, Hf) (Cu, Ni, Ag) Al system, the glass-forming ability of the alloys was improved by the substitution of Co for Ag in the component of late transition metals group. Ó 2004 Elsevier B.V. All rights reserved. PACS: kf; Pf; Fb; Dk 1. Introduction * Corresponding author. Tel.: ; fax: address: jianxu@imr.ac.cn (J. Xu). As a new family of engineering materials, bulk metallic glasses (BMGs) have attracted considerable attention in recent years [1 3]. It has been widely accepted that the presence of multiple components is an important factor to improve the glass-forming ability (GFA) of the alloys [1,4 7]. In other words, making an alloy system more complex is a practical way to search for the new easy glass-forming alloy compositions. In higher-order alloy systems, it is more difficult for the concentrations of all the elements to simultaneously satisfy the composition requirements for the nucleation of crystalline phases than in lower-order systems. This concept of confusing and frustrating crystallization was summed up as the Ôconfusion principleõ [5]. In light of this argument, an equiatomic substitution approach, where the multicomponent glass-forming alloy is consisted of early (ETM) and late (LTM) transition metals with equiatomic composition, was developed in the (Ti, Zr, Hf) (Cu, Ni, Ag) Al system recently by CantorÕs group [8 10]. Indeed, it was revealed that such a multicomponent metallic glass exhibited a wide supercooled liquid region, defined as the temperature interval (DT x ) between the glass transition temperature (T g ) and the onset temperature of crystallization (T x ), DT x = T x T g. The DT x value of (Ti 0.33 Zr 0.33 Hf 0.33 ) 50 (Ni 0.33 Cu 0.33 Ag 0.33 ) 40 Al 10 glasses was reported to be 103 K [8,9]. However, whether such an alloy family is a real BMG former have not been well understood yet. Moreover, InoueÕs group [11] showed that the bulk glassy rod of 1.5mm in diameter was achieved for the Ti 20 Zr 20 Hf 20 Cu 20 Ni 20 alloy. The /$ - see front matter Ó 2004 Elsevier B.V. All rights reserved. doi: /j.jnoncrysol

2 L.C. Zhang, J. Xu / Journal of Non-Crystalline Solids 347 (2004) Ti 20 Zr 20 Hf 20 Cu 20 Ni 20 glass exhibited a high compressive fracture strength of 1920 MPa, which is comparable to the Zr- or Cu-based BMGs [1,3]. Therefore, it is of considerable interest to investigate further the GFA of alloys with the feature of equiatomic substitution. It has been proposed that the GFA of the alloy systems can be improved by a substitution with the elements having a near-zero heat of mixing in the liquid state in order to relax atomic-level stress [12,13]. Thus, for the (Ti, Zr, Hf) (Cu, Ni, Ag) Al system, element Ag in the LTM group is unlikely to be the optimal choice because it has a very positive heat of mixing with Cu or Ni [14]. In the current work, the element Co was chosen to substitute Ag in the (Ti, Zr, Hf) (Cu, Ni, Ag) Al system due to its small heat of mixing with Cu or Ni [14]. Metallic glasses with the nominal compositions of (Ti 0.33 Zr 0.33 Hf 0.33 ) x (Cu 0.33 Ni 0.33 Co 0.33 ) 90 x Al 10 (x = 50, 55, 60), denoted as (TiZrHf) x (CuNiCo) 90 x Al 10 hereafter, were prepared by melt-spinning to investigate the GFA of this series alloys. The content of ETM elements in the system was varied in the range x =50 60at.% due to the fact that most of ETM-based BMG forming alloys have a composition near this range, such as Zr 55 Al 10 Ni 5 Cu 30 [15], Zr 57 Ti 5 Al 10 Cu 20 Ni 8 [16], Zr Nb 2.8 Cu 15.6 Ni 12.8 Al 10.3 [17], Ti 50 Zr 5 Cu 25 Ni 15 Sn 5 [18], Hf 50 Cu 30 Ni 10 Al 10 [19]. 2. Experimental With elemental pieces having the purity higher than 99.9 wt%, master alloys with a nominal composition of (TiZrHf) x (CuNiCo) 90 x Al 10 (x = 50, 55, 60) (in atomic percentage) were prepared by arc melting under a Tigettered argon atmosphere. The metallic glassy ribbons of the alloy were prepared in an argon atmosphere by induction melting the master alloy ingot in a quartz crucible and ejecting it onto a single-roller using a Bühler melt-spinner. The surface speed of the copper roller varies ranging from 5 to 39m/s. The as-quenched ribbons were approximately 4 mm wide and lm thick. The structure of the melt-spun ribbons was analyzed by X-ray diffraction (XRD) using a Rigaku D/max 2400 diffractometer with monochromated Cu Ka radiation (k = nm). Samples for transmission electron microscopy (TEM) were prepared from the 3mm discs by electropolishing in a mixture of ethanol and perchloric acid (9:1 by volume) at 258K with an electrode voltage of 30 90V. The TEM observation was carried out in a JEOL JEM-2000 EX microscope. The glass transition and crystallization of the as-prepared glassy alloys were examined by differential scanning calorimetry (DSC) in a Perkin-Elmer DSC7 under flowing purified argon. Samples with a weight of about 15mg were loaded in alumina pans. A second run under identical conditions was used to determine the baseline after each run, using a heating rate of 20 K/min. To determine the onset temperature of glass transition, T g, the as-prepared samples were first heated to a temperature below the onset temperature of crystallization, T x to reach a relaxed state. The baseline is thus better resolved and the T g can be readily determined from a subsequent run. The melting behavior of the alloys was measured in a Netzsch 404 DSC with alumina container, using a heating rate of 20 K/min. Vickers microhardness tests were performed on a Mitutoya MNK-H3 hardness testing machine at a load of 50g applied for 10s, with about 20 measurements for each sample. 3. Results and discussion Fig. 1 shows the relationship between the ribbon thickness and the wheel speed for the melt-spun (TiZrHf) x (CuNiCo) 90 x Al 10 (x = 50,55,60) alloys. The thickness of the as-quenched ribbons increased as the wheel speed slowed down, ranging from about 30 lm for 39 m/s to 180 lm for 5 m/s. Fig. 2(a) (c) displays the XRD patterns of the as-quenched ribbons prepared with different wheel speeds. As seen in Fig. 2(a), a single amorphous phase formed in the alloy at x = 50 when the wheel speeds were higher than 10m/s, as evidenced by a broad diffuse diffraction maximum at 2h = For the ribbons prepared with a wheel speed of 5 m/s, the formation of crystalline phases such as AlCuHf (S.G.: P6 3 /mmc, a = nm, c = nm) and Ti 2 Cu 3 (S.G.: P4/nmm, a = nm, c = nm) is predominant. The position the crystalline phases in the patterns was determined after separating the amorphous phase and crystalline phases in the line profile using Rietveld method [20]. As a result, the Ribbon thickness (m) x= 50 x= 55 x= Wheel speed (ms -1 ) Fig. 1. Relationship between ribbon thickness and wheel speed for the melt-spun (TiZrHf) x (CuNiCo) 90 x Al 10 alloys.

3 168 L.C. Zhang, J. Xu / Journal of Non-Crystalline Solids 347 (2004) Fig. 3. TEM micrographs of the (TiZrHf) 50 (CuNiCo) 40 Al 10 glassy alloy prepared with a 20m/s of wheel speed: (a) bright-field image, corresponding the selected area electron diffraction patterns (inset) and (b) dark-field image. Fig. 2. XRD patterns of the melt-spun (TiZrHf) x (CuNiCo) 90 x Al 10 alloys prepared with different wheel speeds: (a) x = 50; (b) x = 55 and (c) x = 60. maximum thickness of fully amorphous ribbon samples for this alloy was determined to be around 85lm. Fig. 3 shows the TEM bright-field micrograph with corresponding selected area electron diffraction (SAED) patterns and dark-field micrograph of the as-quenched alloys at x = 50 prepared with the wheel speed of 20 m/s. Only the featureless amorphous phase was observed without any quenched-in nuclei or residual crystallites, as seen in Fig. 3(a) and (b). The broad halo in the SAED pattern indicates the formation of the uniformly amorphous phase in the as-quenched state. Similarly, crystalline phases are detectable in the XRD patterns for the ribbons when the wheel speed is lower than 10m/s for the alloy at x = 55 or 20m/s for x = 60, as seen in Fig. 2(b) and (c). Comparing the fully glassy samples of the three alloys, it can be seen that, with increasing the content of ETM elements, the diffuse diffraction maximum shifted to a lower 2h angle side. The wavenumber Q p defined as Q p =4psinh/k, of broad diffuse maximum of the glassy alloys decreased from 27.24nm 1 for x =50 to 26.56nm 1 for x = 60. In addition, the fully glassy ribbons of all the three alloys are ductile and can be bend elastically over 180 without fracture. The microhardness of the glassy alloys are insensitive to the change in composition, as measured to be around 6200 ± 350 MPa. As such, the yield strength, corresponding to typically 1/3 of the microhardness, can be estimated to be about 2.1GPa. Fig. 4(a) (c) present the DSC scans of the asquenched ribbons prepared with the various wheel speeds for the (TiZrHf) x (CuNiCo) 90 x Al 10 alloys, in

4 L.C. Zhang, J. Xu / Journal of Non-Crystalline Solids 347 (2004) m/s 10 m/s 20 m/s 30 m/s 0.1 W/g m/s 10 m/s 20 m/s 30 m/s 0.1 W/g m/s 20 m/s 30 m/s 39 m/s 0.1 W/g which the dash line curves are for the relaxed samples by a preheating run. As seen in Fig. 4(a), for the samples of the fully glassy alloy at x = 50 formed using a wheel speed higher than 10 m/s, two main exothermic peaks due to crystallization of the amorphous phase appear, with the peak temperatures around 800 and 870 K, respectively. It was observed that before the onset of the first crystallization peak, the glass transition, seen as an endothermic signal, appears at the temperature T g T g Fig. 4. DSC scans of the melt-spun (TiZrHf) x (CuNiCo) 90 x Al 10 alloys prepared with different wheel speeds: (a) x = 50; (b) x = 55 and (c) x = 60 (with a heating rate of 20K/min, the dash line curves are for the relaxed samples by a preheating run). (a) (b) (c) about 705 K. Additionally, a broad exothermic reaction in a range from 650 to 750K caused by an irreversible relaxation is observed. The heat release caused by relaxation increased with increasing wheel speed. For the relaxed sample after a preheating run, the glass transition temperature, T g and the onset temperature of crystallization, T x1 were determined to be 657 and 781K, respectively. Then, the width of supercooled liquid region, DT x was obtained to be 124 K. By integrating the peak areas from the exothermic peaks, the enthalpy of crystallization for the two steps, DH x1 and DH x2 is 1.07 and 1.24 kj/mol, respectively. Moreover, even for the sample obtained with a wheel speed of 5m/s, a small exothermic peak associated with the second crystallization is still visible. It means that a small fraction of the amorphous phase remained in the ribbons with a thickness of 180 lm. Similar to the results of the alloy at x = 50, two-steps crystallization also happens in the fully glassy alloy of x = 55 and 60 prepared with the wheel speed higher than 10m/s, as seen in Fig. 4(b) and (c). The glass transition signal is detectable also in the alloy at x = 55, but not at x = 60. The thermal properties obtained with DSC measurement for the three glassy alloys are listed in Table 1, including the T g, the onset crystallization temperature and enthalpy of crystallization at two steps, T x1, T x2, DH x1 and DH x2, the width of supercooled liquid region, DT x and the total heat release of crystallization, DH total defined as the sum of H x1 and DH x2. Comparing the samples of fully glassy alloys, the enthalpy of crystallization at two steps, especially for the first crystallization, was reduced for the samples formed at a wheel speed of 10 m/s in both alloys. It implies that partially crystallization occurred in the as-quenched ribbons for the alloys of x = 55 and 60. Fully glassy alloys is not achievable under this condition, which is in agreement with the XRD findings as shown in Fig. 2(b) and (c). Therefore, fully glassy ribbons can be obtained only under the condition that the wheel speed exceeds 20m/s for the alloy at x = 55 and 30 m/s for x = 60. The maximum thickness of fully glass formation for the alloy of x =55 and 60 was determined to be around 55 and 40 lm, respectively. According to the above results from XRD and DSC, the maximum thickness of fully glassy ribbon varies with the content of the ETM elements (Ti, Zr, Hf). This dependence is plotted in Fig. 5. It is indicated that with increasing the content of ETM element, the maximum thickness of glass formation reduced. In other words, the GFA of the alloys was degraded. Among the present three alloys, the best glass former is the (TiZrHf) 50 (CuNiCo) 40 Al 10 alloy. To identify the structural changes related to the two main crystallization events in the DSC scans, the samples of fully glassy alloys were continuously heated in the DSC to the end temperatures of the two exothermal

5 170 L.C. Zhang, J. Xu / Journal of Non-Crystalline Solids 347 (2004) Table 1 Thermal properties obtained from the DSC measurements for the melt-spun (TiZrHf) x (CuNiCo) 90 x Al 10 glassy alloys Alloys T g (K) T x1 (K) T x2 (K) DT x (K) DH x1 (kj/mol) DH x2 (kj/mol) DH total (kj/mol) x = ± ± ± ± ± ± ± 0.22 x = ± ± ± 1 97 ± ± ± ± 0.20 x = 60 N/A 730 ± ± 1 N/A 0.92 ± ± ± 0.24 Ribbon thickness (µm) Am. Amorphous Am.+Cyst. Am. + Cryst (TiZrHf) content (at.%) Fig. 5. Map of phase formation for the (TiZrHf) x (CuNiCo) 90 x Al 10 alloys with various ribbon thickness. events, as marked by dots in the curves, as seen in Fig. 4, and then cooled at 320K/min to room temperature, for XRD measurements. The XRD patterns are shown in Fig. 6. As seen in Fig. 6(a), at the end of the first-stage crystallization corresponding to the first exothermal event, no obvious crystalline diffraction peaks appeared. It is probably due to a reason that the size of crystallized phase is too small, in the nanometer scale, to be identified from the patterns. Similar results were observed also in many multicomponent metallic glasses that the primary phase precipitated as the nanocrystals after primary crystallization performed [10,21,22]. It can be seen in Fig. 6(b) that, at the second stage of crystallization, the crystallized products of the three alloys are the same, consisting of the intermetallic phases of AlCuHf and Ti 2 Cu 3. This result indicates that the composition change in the current range for this system has only a minor effect on the final crystallized phases. Furthermore, to determine the reduced glass transition temperature, T rg defined as the glass transition temperature T g divided by the liquidus temperature T l [23], i.e., T rg = T g /T l which is normally used as a indicator to evaluate the GFA of the alloys, the melting behavior of the three alloys was investigated. The DSC curves at a temperature range near the melting points for the three alloys are shown in Fig. 7(a). In all cases, the melting event exhibited several peaks. These DSC features suggest that the three alloys of (TiZrHf) x (CuNiCo) 90 x Al 10 (x = 50, 55, 60) are far from a eutectic composition. The (a) x=50 T l x=55 T m x=60 (b) Fig. 6. XRD patterns of the (TiZrHf) x (CuNiCo) 90 x Al 10 glassy alloys heated beyond (a) the first and (b) the second exothermic peaks at 20K/min, then followed by cooling to room temperature in DSC. Fig. 7. DSC curves in a range near the melting temperature for (a) the (TiZrHf) x (CuNiCo) 90 x Al 10 alloys and (b) (TiZrHf) 50 (CuNiAg) 40 Al 10 (with a heating rate of 20K/min).

6 L.C. Zhang, J. Xu / Journal of Non-Crystalline Solids 347 (2004) melting temperature T m and liquidus temperature T l of the three alloys, obtained from the DSC curves, depends on the content of the ETM elements (Ti, Zr, Hf), as shown in Fig. 8. It was indicated that with increasing the content of ETM elements group, both T m and T l dropped, but the temperature interval (from the onset to the end) spanned by the melting process was extended from 94K for the alloy at x = 50 up to 188K at x = 60.A better GFA of the alloys seems to associate with a narrower temperature span. The T rg of the alloys at x = 50 and 55 was calculated to be and 0.468, respectively. Apparently, the values for this alloy system are much smaller than those of bulk glass-forming alloys, such as 0.55 for Mg 65 Cu 25 Y 10, 0.59 for Zr 57 Ti 5 - Al 10 Cu 20 Ni 8, 0.64 for Zr 41.2 Ti 13.8 Cu 12.5 Ni 10 Be 22.5, and 0.69 for Pd 40 Cu 30 Ni 10 P 20, as summarized in Ref. [24]. This implies that the GFA of the (TiZrHf) (CuNiCo) Al alloys based on the equiatomic substitution strategy is not so strong to compare with the Mg-, Zr- or Pdbased bulk glasses, even though the system contains more number of the components. The investigation on whether the BMG former exists in the (TiZrHf) (CuNiCo) Al system is underway. For comparison, the XRD patterns of the melt-spun (TiZrHf) 50 (CuNiAg) 40 Al 10 alloy prepared with two wheel speeds, 30 and 39m/s, are shown in Fig. 9. It can be seen that the fully glassy ribbon was obtained only in the case using the wheel speed higher than 39m/s. In the pattern of the as-quenched ribbon obtained using the wheel speeds of 30 m/s, the metastable fcc NiZr 2 (big cube) phase is detectable, overlapped with the broad diffuse diffraction of the glassy phase, indicating the fully glassy ribbon is not achievable. Evidently, the GFA of this alloy is weaker than that of any alloy among the three Co-substituted alloys. In other words, 1500 solid: Ag-containing Fig. 9. XRD patterns of the melt-spun (TiZrHf) 50 (CuNiAg) 40 Al 10 alloy prepared with two wheel speeds. the (TiZrHf) (CuNiCo) Al system has the GFA significantly better than (TiZrHf) (CuNiAg) Al system. The crystallization of the (TiZrHf) 50 (CuNiAg) 40 Al 10 glassy alloy has been investigated in the previous work [25]. The T g and DT x of the glassy alloy were given to be 680 and 53 K, respectively. The DSC curve associated with the melting behavior is displayed in Fig. 7(b). Besides the main peak, the melting proceeds through several endothermic events shown as the small peaks, implying that the alloys located at a eutectic-off composition. The T m and T l of the alloy were determined to be 1197 and 1413 K, respectively, as illustrated by solid symbol in Fig. 8. Then, the T rg of this alloy is given as 0.481, which is comparable to the value of the Co-substituted alloys. It indicates that the T rg is not a sensitive indicator to evaluate the GFA of these alloys. In addition, the Ag-containing alloy exhibits a wide temperature span between T m and T l about 216 K. It is a possible source of the findings that GFA of this alloy is not so stronger as the Co-substituted alloys (TiZrHf) Content (at.%) Fig. 8. Dependence of the melting temperature T m and liquidus temperature T l on the (Ti, Zr, Hf) content in the (TiZrHf) x (CuNiCo) 90 x Al 10 alloy, with the data of the (TiZrHf) 50 (CuNiAg) 40 Al 10 alloy as marked by solid symbols for comparison. T l T m 4. Conclusion For the (TiZrHf) x (CuNiCo) 90 x Al 10 (x = 50, 55, 60) multicomponent alloys with the equiatomic substitution, fully amorphous ribbons can be obtained by melt spinning. By decreasing the content of the ETM (Ti, Zr, Hf) elements, the glass-forming ability of the alloys was improved. The maximum thickness of fully glassy ribbons was increased from 40 lm for the alloy at x =60to85lm for the alloy at x = 50. The (TiZrHf) 50 (CuNiCo) 40 Al 10 glassy alloy exhibits the largest supercooled liquid region of 124 K, and a reduced glass transition temperature of The substitution of Co for Ag in the (TiZrHf) (CuNiAg) Al alloy system significantly enhanced the glass-forming ability of the alloys.

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