Enhanced Amorphization by Sn Substitution for Si and B in the Ball-milled Ti 50 Ni 22 Cu 18 Al 4 Si 4 B 2 Alloy. L.C. Zhang, Z.Q. Shen, J.

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1 Journal of Metastable and Nanocrystalline Materials Vols (2004) pp online at Journal 2004 Citation Trans Tech (to Publications, be inserted by Switzerland the publisher) Copyright by Trans Tech Publications Enhanced Amorphization by Sn Substitution for Si and B in the Ball-milled Ti 50 Ni 22 Cu 18 Al 4 Si 4 B 2 Alloy L.C. Zhang, Z.Q. Shen, J. Xu * Shenyang National Laboratory for Materials Science, Institute of Metal Research, CAS, 72 Wenhua Road, Shenyang, , China Keywords: Mechanical Alloying, Amorphous Alloys, Supercooled Liquid Region, Titanium Abstract Element Sn is substituted for metalloid elements Si and B in the Ti 50 Ni 22 Cu 18 Al 4 Si 4 B 2 alloy to form the Ti 50 Ni 22 Cu 18 Al 4 Sn x (Si 0.67 B 0.33 ) 6-x (x=0, 3, 6) series alloys for high-energy ball milling. For comparison, the amorphous alloy ribbon with a composition of x=6 was prepared using melt-spinning method as well. Structural features of the samples were characterized using x-ray diffraction, transmission electron microscopy and differential scanning calorimetry. In the ball-milled alloys, the fraction of residual α-ti crystallites in the final products was reduced with increasing the Sn content. For the Ti 50 Ni 22 Cu 18 Al 4 Sn 6 alloy, without any metalloid element, the complete amorphization is achievable. Such a ball-milled glassy alloy shows the thermal properties comparable with those of the melt-spun glassy ribbon with the same composition, and exhibits the largest supercooled liquid region about 66 K. 1. Introduction It has been revealed that the Ti-based amorphous alloys have ultrahigh tensile fracture strengths (σ f ) around 2 GPa [1-3]. They are promising as a new family of light-weight materials with high specific strength. Unfortunately, it has been difficult to prepare the Ti-based metallic glasses in bulk form. The largest thickness achieved so far for Ti-based multicomponent amorphous alloys using conventional casting methods is only of the order of 1 mm [3-6]. Consequently, the powder processing and subsequent warm consolidation in the vicinity of glass transition temperature of the glassy alloys is still an attractive approach to prepare the truly bulky materials [7-13]. For this feature, it is an important issue to synthesize the glassy alloy powders with a sizable supercooled liquid region, defined as the temperature interval ( T x ) between the glass transition temperature (T g ) and the onset temperature of crystallization (T x ). Recently, it has been found that the Ti-based glassy alloys with the T x value about 60 K can be prepared by mechanical alloying (MA) of the Ti-Cu-Ni-Si-B multicomponent alloy system [14, 15]. However, the MA-induced amorphization is incomplete, where the nanoscale particles of unreacted α-ti phase with a volume fraction of approximately 20% remained in the final ball-milled products [14]. It was assumed that a factor that the system contains Si and B elements is possible source responsible for the presence of residual crystals in the resulting powders. As usual, the metalloid elements lack of plasticity are uneasy to be alloyed under ball-milling processing. In the present work, the Ti 50 Ni 22 Cu 18 Al 4 Si 4 B 2 alloy, which yielded a large T x about 64 K in the MA glassy powders [14], was chosen as the base composition. Metalloid elements, Si and B, are replaced with ductile metal Sn to form the alloys of Ti 50 Ni 22 Cu 18 Al 4 Sn x (Si 0.67 B 0.33 ) 6-x (x=0, 3 and 6). The additional consideration to introduce the Sn is due to the finding that the minor addition of Sn improved the thermal stability of the supercooled liquid in the Ti-based glasses [3, 16-18]. The effect of Sn substitution on the milling-induced amorphization and crystallization of the milled glassy alloys was investigated. * Address all correspondence to this author. address: jianxu@imr.ac.cn

2 Journal of Metastable and Nanocrystalline Materials Vols Title of Publication (to be inserted by the publisher) 2. Experimental With elemental pieces having the purity higher than 99.9 wt.%, the master alloys with a nominal composition of Ti 50 Ni 22 Cu 18 Al 4 Sn x (Si 0.67 B 0.33 ) 6-x (x=0, 3 and 6; in atomic percentage) were prepared by arc melting under a Ti-gettered argon atmosphere. The alloyed buttons were then crushed into fragments that were used as starting materials for ball milling. The prealloyed fragments together with hardened steel balls were loaded in a hardened steel vial under an argon-filled glove box with less than 1 ppm O 2 and H 2 O. A ball-to-powder weight ratio about 5:1 was employed. The ball milling process was performed in a SPEX 8000 shaker mill cooled by forced flowing air. For all the samples discussed in this paper, the milling time was fixed at 32 hrs. This time was chosen based on the results of our prior studies of this system [19], to assure that the milling has proceeded sufficiently long for the samples to reach a steady state, but short enough to avoid excessive contamination from the milling media. The metallic glassy ribbons of the alloy at x=6 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 was 39 m/s. The as-quenched ribbons were approximately 4 mm wide and µm thick. The amorphicity of the MA powders and melt-spun (MS) ribbons were analyzed by x-ray diffraction (XRD) using a Rigaku D/max 2400 diffractometer with monochromated Cu K α radiation (λ= nm). For the MA powders, samples for transmission electron microscopy (TEM) were prepared by first embedding the powder particles in a nickel foil via electrodeposition for mechanically thinning and subsequently ion-milling to electron transparency. The TEM observation was carried out in a JEOL JEM-2010 microscope. The glass transition and crystallization of the glassy alloys formed were examined by differential scanning calorimetry (DSC) in a Perkin-Elmer DSC7 under flowing purified argon. Isothermal measurement was carried out by heating up to the desired temperature at a heating rate of 200 K/min. The iron and oxygen contents in the ball-milled products were examined using inductively coupled plasma emission spectroscopy (ICP10P, ARL) and a LECO TC-436 system to be 0.50 wt. % and 0.15 wt.%, respectively. 3. Results and Discussion Figure 1 presents the XRD patterns of the prealloyed fragments with various compositions Ti 50 Ni 22 Cu 18 Al 4 Sn x (Si 0.67 B 0.33 ) 6-x (x=0, 3 and 6) after milling for 32 hrs. For comparison, a pattern of the MS ribbon for the alloy at x=6 is plotted together. In all these patterns, a broad diffuse diffraction maximum at 2θ=35 55 indicates the formation of amorphous phase in the MA powders. However, crystalline peaks of α-ti phase are still identifiable in the pattern of the Sn-free alloy even though the milling time has been extended to a steady state for the Fig. 1 XRD patterns of the ball-milled Ti 50 Ni 22 Cu 18 Al 4 Sn x (Si 0.67 B 0.33 ) 6-x alloys.

3 490 Metastable, Mechanically Alloyed and Nanocrystalline Materials 2003 structural evolution. It was noted that as the Sn content increased, the diffraction peaks of α-ti crystallites disappear in the patterns. It suggests that the MA-induced amorphization in the Sn-containing alloys is more complete than in the Sn-free alloy. For the alloy at x=6, the wavenumber Qp, defined as Qp=4πsinθ/λ, of broad diffuse maximum for the ball-milled glassy alloy was determined to be nm-1, which is well in agreement with the value, Qp=26.67 nm-1, for the MS alloy. Also, it reflects that the amorphous phase obtained using different preparation methods is very similar. Figure 2 (a)-(c) illustrates the TEM bright-field micrographs and corresponding selected area electron diffraction (SAED) patterns of the MA Ti50Ni22Cu18Al4Snx(Si0.67B0.33)6-x powders. For the MA alloy without Sn, the α-ti particles of about 10 nm in size, with an areal fraction of approximately 20%, dispersed in the featureless amorphous matrix, as shown in Fig. 2 (a). With the 3%Sn substitution, the fraction of the residual α-ti particles was significantly reduced in the milled products, as seen in Fig. 2 (b). Up to x=6, the final milled product is a single amorphous phase whereas the unreacted crystallites are no longer visible, as seen in Fig. 2 (c). The selected area electron diffraction (SAED) pattern shows only a broad halo. As a result, the TEM observation further confirms that the complete glass formation under high-energy ball milling can be achieved only in the Ti50Ni22Cu18Al4Sn6 alloy. Namely, it can be concluded that the Sn substitution for metalloid elements is effective to enhance the amorphization in the MA Ti-based multicomponent alloys. Fig. 2 TEM bright-field micrographs and corresponding selected area electron diffraction patterns (inset) of the ball-milled Ti50Ni22Cu18Al4Snx(Si0.67B0.33)6-x alloys: a) x=0, b) x=3 and c) x=6. Figure 3 displays the DSC scans in the mode of continuous heating with a heating rate of 40 K/min for the samples with different Sn contents, together with a scan of the MS glassy alloy at x=6. In all cases, an endothermic signal associated with the glass transition and an exothermic reaction due to crystallization is evident. The crystallization of the glassy alloys was performed through a single step. With increasing Sn substitution, the onset of glass transition temperature, Tg, and the onset temperature of crystallization, Tx, occurs at lower temperatures, accompanied by somewhat extension of the supercooled liquid region. The largest Tx value, 66 K, was obtained at the x=6 alloy. The heat

4 Journal of Metastable and Nanocrystalline Materials Vols Title of Publication (to be inserted by the publisher) Fig. 3 DSC scans of the ball-milled Ti 50 Ni 22 Cu 18 Al 4 Sn x (Si 0.67 B 0.33 ) 6-x alloys. Fig. 4 XRD patterns of the ball-milled Ti 50 Ni 22 Cu 18 Al 4 Sn x (Si 0.67 B 0.33 ) 6-x alloys after heating to the temperature beyond the exothermic event in DSC. of crystallization obtained by integrating the area under the peak, H x, is listed in Table 1, together with the T g, T x, and T x. For comparison purpose, the corresponding data for the MS glassy alloy at x=6 are included in Table 1 as well. With increasing the Sn content, the heat of crystallization increased from 1.67 kj/mol for the Sn-free alloy to 2.68 kj/mol for the 6%Sn alloy, indicating that the volume fraction of the glassy phase in the milled product was significantly increased. For the 6%Sn alloy, the T x and H x of the MA glassy alloy is comparable with those of the MS glass, but the T g and T x of the MA glassy alloy is 16 K and 18 K lower than that of the MS glass, respectively. Even so, the XRD results and the thermal properties obtained from DSC measurement show that the glassy phase formed through the two preparation methods is very similar. Similar result was observed also in the Zr-Al-Ni-Cu-Co [20] and Ti-Zr-Hf-Cu-Ni-Ag-Al [21] multicomponent glassy alloys prepared by both methods. Table 1 Glass transition temperature (T g ), onset temperature of crystallization (T x ), extension of the supercooled liquid region (ΔT x ) and heat of crystallization (ΔH x ) obtained from DSC measurements for the Ti 50 Ni 22 Cu 18 Al 4 Sn x (Si 0.67 B 0.33 ) 6-x glassy alloys. Alloys Synthesis route T g, K T x, K ΔT x, K ΔH x, kj/mol x=0 MA ±0.02 x=3 MA ±0.12 x=6 MA ±0.14 x=6 MS ±0.01

5 492 Metastable, Mechanically Alloyed and Nanocrystalline Materials 2003 The XRD patterns in Fig. 4 identify the structural changes associated with the exothermic event in the DSC traces. The samples were continuously heated in the DSC to a temperature at which the exothermic reaction is completed for each alloy, as marked by dots in the curves in Fig. 3, then cooled at 320 K/min to room temperature for XRD measurements. For all three alloys, the XRD patterns showed that the amorphous phase transformed into the cubic NiTi-type phase and an unidentified phase, in a way of eutectic-type crystallization [22]. In addition, the crystallized phases for the ball-milled alloys is the same as those for the MS glass of x=6. Evidently, the resulting crystallized phases are irrespective of the Sn substitution for Si and B and the presence of unreacted crystallites in the samples. Isothermal DSC traces for the three MA glassy alloys are shown in Fig. 5. The annealing temperature was below the T x obtained by continuous heating for each alloy. For all cases, one distinct exothermic signal after a certain incubation period was displayed. The form of the heat flow signals is typical shape for a nucleation and growth process [23]. It proves that the MA products of the alloys formed the amorphous phases rather than the mixtures of nanograined intermetallics. 4. Summary The substitution of element Sn for metalloid elements Si and B in the Ti 50 Ni 22 Cu 18 Al 4 Si 4 B 2 multicomponent alloy significantly enhanced the glass formation under high-energy ball milling. With increasing the Sn content, the fraction of residual α-ti crystallites was reduced. For the Ti 50 Ni 22 Cu 18 Al 4 Sn 6 alloy, without any Si and B, the complete amorphization is achievable in the final ball-milled products. Such a glassy alloy shows the thermal properties quite similar to those of the melt-spun glassy ribbon with the same composition. A wide supercooled liquid region about 66 K is obtained. The Sn substitution has minor effect on the crystallization behavior of the glassy alloys. The glassy phases formed by MA crystallize through the nucleation and growth of intermetallics. Acknowledgements The authors gratefully acknowledge the financial support from the National Natural Science Foundation of China under contracts No and J.X. is indebted to the K.C. Wong Education Foundation, Hong Kong, for supporting this project. References Fig. 5 Isothermal DSC scans at several temperatures for the ball-milled Ti 50 Ni 22 Cu 18 Al 4 Sn x (Si 0.67 B 0.33 ) 6-x alloys. [1] A. Inoue, H.M. Kimura, T. Masumoto, C. Suryanarayana and A. Hoshi: J. Appl. Phys. Vol. 51 (1980), p [2] A. Inoue, K. Amiya, A. Katsuya and T. Masumoto: Mater. Trans. JIM Vol. 36 (1995), p. 858 [3] T. Zhang and A. Inoue: Mater. Sci. Eng. Vol. A (2001), p. 771 [4] T. Zhang and A. Inoue: Mater. Trans. JIM Vol. 40 (1999), p. 301 [5] Y.C. Kim, W.T. Kim and D.H. Kim: Mater. Trans. Vol. 43 (2002), p [6] G. He, W. Löser, J. Eckert and L. Schultz: J. Mater. Res. Vol. 17 (2002), p [7] Y. Kawamura, H. Kato, A. Inoue and T. Masumoto: Appl. Phys. Lett. Vol. 67 (1995), p. 2008

6 Journal of Metastable and Nanocrystalline Materials Vols Title of Publication (to be inserted by the publisher) [8] Y. Kawamura, H. Kato, A. Inoue and T.Masumoto: Int. J. Powder. Metall. Vol 33 (1997), p. 50 [9] D.J. Sordelet, E. Rozhkova, P. Huang, P.B. Wheelock, M.F. Besser, M.J. Kramer, M. Calvo-Dahlborg and U. Dahlborg: J. Mater. Res. Vol. 17 (2002), p. 186 [10] T. Itoi, T. Takamizawa, Y. Kawamura and A. Inoue: Scripta Mater. Vol. 45 (2001), p [11] S. Ishihara, W. Zhang and A.Inoue: Scripta Mater. Vol. 47 (2002), p. 231 [12] M.H. Lee, D.H. Bae, W.T. Kim, D.H. Kim, E. Rozhkova, P.B. Wheelock and D.J. Sordelet: J. Non-Cryst. Solids Vol. 315 (2003), p. 89 [13] J. Robertson, J.T. Im, I. Karaman, K.T. Hartwig and I.E. Anderson: J. Non-Cryst. Solids. Vol. 317 (2003), p [14] L.C. Zhang, J. Xu and E.Ma: J. Mater. Res. Vol. 17 (2002), p [15] I.-K. Jeng, P.-Y. Lee, J.-S. Chen, R.-R. Jeng, C.-H. Yeh and C.-K. Lin: Intermetallics Vol. 10 (2002), p [16] T. Zhang and A. Inoue: Mater. Trans. JIM Vol. 39 (1998), p [17] Y.C. Kim, S. Yi, W.T. Kim and D.H. Kim: Mater. Sci. Forum Vol (2001), p. 67 [18] Y.C. Kim, S. Yi, W.T. Kim and D.H. Kim: Mat. Res. Soc. Symp. Proc. Vol. 644 (2001), p. L4.9.1 [19] L.C. Zhang and J. Xu: Mater. Sci. Forum Vol (2002), p. 47 [20] A. Sagel, R.K. Wunderlich, J.H. Perepezko and H.J. Fecht: Appl. Phys. Lett. Vol. 70 (1997), p. 580 [21] L.C. Zhang, Z.Q.Shen and J. Xu: J. Mater. Res. Vol. 18 (2003), p [22] U. Köster: Z. Metallkde. Vol. 75 (1984), p. 691 [23] L.C. Chen and F. Spaepen: J. Appl. Phys. Vol. 69 (1991), p. 679