Compression stress induced flow temperature reduction in a bulk Zr 41:2 Ti 13:8 Cu 12:5 Ni 10:0 Be 22:5 metallic glass

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1 Scripta Materialia 47 (2002) Compression stress induced flow temperature reduction in a bulk Zr 41:2 Ti 13:8 Cu 12:5 Ni 10:0 Be 22:5 metallic glass H.J. Jin, X.J. Gu, F. Zhou, K. Lu * Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang , China Received 28 December 2001; received in revised form 22 January 2002; accepted 25 June 2002 Abstract For a bulk metallic glass Zr 41:2 Ti 13:8 Cu 12:5 Ni 10:0 Be 22:5, flow behavior was investigated under uniaxial pressure up to 1.37 GPa upon continuous heating. Under compression stress, flow temperature is reduced by 45 K/1.37 GPa and the processible temperature range is expected to be possibly widened by about 70 K/l.37 GPa. Ó 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Metallic glasses; Hot pressing; Plastic; Stress; Flow 1. Introduction Metallic glasses usually exhibit significant plasticity in the supercooled liquid region (DT x ¼ T x T g, where T x is the crystallization temperature and T g the glass transition temperature), due to a drastic drop in viscosity (usually by several orders of magnitude) during glass transition upon heating [1]. Accordingly, it is possible to make net-shape products by applying superplastic deformation in this temperature region [2,3]. It is also possible to prepare full density bulk metallic glasses with a dimension independent of the critical cooling rate by consolidating amorphous powders or ribbons in the supercooled liquid region [4,5]. However, most metallic glass formers exhibit a narrow, even no supercooled liquid region upon heating. In * Corresponding author. Fax: address: lu@imr.ac.cn (K. Lu). some studies, external forces (e.g. pressure) have been attempted to retard the crystallization in order to enlarge DT x [4,6]. But unexpectedly, for some metallic glasses an applied pressure can enhance crystallization and reduce DT x [7]. Therefore, it is expected to carry out the plastic deformation or consolidation at a temperature as low as possible. Thus it is necessary to study the flow behavior around and below T g in metallic glasses. It was found that the metallic glasses could be deformed homogeneously at temperatures above 0:7T g [8]. Several recent studies have focused on the study of flow behavior, and its temperature and strain rate dependence for bulk metallic glasses [9,10]. However, most of previous studies were carried out by measuring the strain stress curves at given strain rates and temperatures. When a metallic glass is held isothermally at temperatures below T g, structural relaxation, even crystallization, usually occurs which leads to the change in its /02/$ - see front matter Ó 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S (02)

2 788 H.J. Jin et al. / Scripta Materialia 47 (2002) properties (including the flow behavior), especially for the measurements performed at temperatures close to T g or last for a long time. Recently, viscous flow behavior of several metallic glasses was measured by thermomechanical analyzer (TMA) under continuous heating. But the applied stress was very low (<50 kpa), and the flow could not be detected clearly below T g [11,12]. In this paper, we will study the flow behavior of a bulk metallic glass Zr 41:2 Ti 13:8 Cu 12:5 Ni 10:0 Be 22:5 upon continuous heating under compression stress up to 1.37 GPa. 2. Experimental A 7 mm thick amorphous Zr 41:2 Ti 13:8 Cu 12:5 - Ni 10:0 Be 22:5 plate was prepared by arc-melting the mixture of pure elements in an argon atmosphere and casting into a copper mold. The amorphism of the specimens was confirmed by means of X-ray diffraction (XRD). The thermal stability of the as-quenched amorphous alloy samples was characterized by means of differential scanning calorimeter (DSC Pyris 1, Perkin Elmer) in a flowing argon atmosphere. The temperature (with an accuracy of 0.02 K) and heat flow of the DSC were calibrated by using pure In and Zn standard samples. Al pans were used for both the sample and the reference. The measurements of flow temperature were carried out on a piston-cylinder pressure vessel with the internal diameter of 8 mm in a furnace with the vacuum of Pa. A thermocouple was placed right under the specimen to monitor the temperature, and the difference between the actual temperature and the measured is less than 3 K. Specimens with a dimension of 5 6 2mm 3 were cut from the same plate. Each specimen was heated up to 653 K that is above the glass transition temperature, then cooled down with the same cooling rate of 5 K/min to ensure the same thermal history for all specimens. Both sides of each specimen were cut to be parallel and were mechanically polished to insure the flatness. The heating was performed at the rate of 5 K/min for all specimens. The data of temperatures and pressures were collected simultaneously upon heating. 3. Results and discussion In the measurements, a preset uniaxial pressure (compression stress) was applied to the sample and held before heating up the sample at a constant rate. With an increase of temperature (T ), the pressure (P) applied to the sample increases due to the thermal expansion of the sample and the piston, as shown in Fig. 1(a). The pressure increase deviates from the linear P T line at a critical temperature (about 600 K in Fig. 1(a)) and then the pressure drops drastically. When temperature exceeds 670 K, pressure increases with temperature again. This is due to the fact that the specimen has been deformed so that it fulfilled the cylinder vessel. The resultant specimen at this temperature was a perfect cylinder with the same diameter as that of the vessel. When heating the sample (a cylinder-shape) directly under the same measurement conditions for the second run, the pressure increases linearly in the whole temperature range as only the thermal expansion was involved during heating. Fig. 1(b) compares the P T curve at a low preset pressure with a DSC trace performed at the same heating rate. It is obvious that the abrupt drop in pressure corresponds to the glass transition process. The pressure drop originates from the sample thickness reduction caused by a homogeneous deformation of the sample occurred in this temperature range. The flow temperature T f, as defined in Fig. 1(b), is very close to the end temperature of glass transition (T g-end ). Fig. 1(c) shows the P T curves performed under various preset uniaxial pressures (compression stresses). An evident decrease of flow temperature with increasing preset pressure was observed. It is also shown that the specimens were compressed into a cylinder in the same temperature segment of about 50 K (in about 10 min) in all cases. The average strain rate was estimated to be s 1 for all specimens, even though the late stage of the deformation was restricted by the inner surface of the vessel. From results shown in Fig. 1(c), the variation of processible temperature range with the applied pressure can be derived for the Zr 41:2 Ti 13:8 Cu 12:5 - Ni 10:0 Be 22:5 amorphous alloy, as shown in Fig. 2.

3 H.J. Jin et al. / Scripta Materialia 47 (2002) Fig. 1. Typical P T curve (a), the DSC trace and the P T curve performed under low stress (b) and P T curves measured with different preset applying stresses (c) for the Zr 41:2 Ti 13:8 Cu 12:5 Ni 10:0 Be 22:5 metallic glass. The heating rate is 5 K/min. For comparison, the onset (T g-onset ) and the end (T g-end ) temperatures of glass transition determined by DSC are also included in Fig. 2. Unlike the supercooled region, the processible temperature range is defined as the temperature range between the flow temperature (T f ) and the crystallization temperature (T x ). It can be clearly seen in Fig. 2 that the flow temperature was reduced by about 45 K/1.37 GPa under compression stress. Since the Fig. 2. Variation of flow temperature with increasing uniaxial pressure and the crystallization temperature with increasing hydrostatic pressure [13]. The processible temperature range was enlarged by about 70 K/l.37 GPa. specimen tends to flow under compression stress at temperatures above T g, only the hydrostatic pressure can be imposed on the supercooled liquid. Therefore, effects of pressure on crystallization behavior are commonly studied under hydrostatic condition in amorphous alloys. Jiang et al. [13] reported that the crystallization temperature increases with the applied pressure by 19 K/GPa for the Zr 41:2 Ti 13:8 Cu 12:5 Ni 10:0 Be 22:5 amorphous alloy (dashed line). Therefore, regardless of the difference in pressures applied in different temperature region, the processible temperature range was enlarged by about 70 K/1.37 GPa for the Zr 41:2 Ti 13:8 Cu 12:5 Ni 10:0 Be 22:5 amorphous alloy at a heating rate of 5 K/min. Temperature dependence of flow stress (stress corresponding to T f ) is summarized in Fig. 3. It is evident that the flow stress begins to drop below calorimetric glass transition temperature and decreases to almost zero around the end of glass transition. Similar experiments have been performed under compressive stresses less than 20 kpa for Pd 40 Cu 30 Ni 10 P 20 amorphous alloy [12]. The length reduction was reported to begin at temperature 48 K higher than T g determined by DSC at the same heating rate of 20 K/min. Since the width of the glass transition region usually

4 790 H.J. Jin et al. / Scripta Materialia 47 (2002) Fig. 3. Flow stress plotted as a function of temperature for the Zr 41:2 Ti 13:8 Cu 12:5 Ni 10:0 Be 22:5 metallic glass. For comparison, room temperature yield stress is also indicated in this figure. increases with the enhanced heating rate [14], the flow is estimated to occur around the end temperature of glass transition, which is in agreement with this study. The room temperature yield stress measured from compression tests for the same amorphous alloy [15] is also indicated in Fig. 3. It should be pointed out that the length to diameter (l=d) aspect ratio of specimens used in compressing tests in Ref. [15] was approximately 3:2, and that of our present specimen is about 2:5.5. Usually, the measured yield stress increases with the decreasing aspect ratio of the specimen. So for the specimens used in our experiments, the room temperature yield stress should be a little higher than that measured in Ref. [15]. It can be concluded from Fig. 3 that at temperatures below T g-onset (620 K), the flow stress (1.37 GPa at 603 K) is comparable to the room temperature yield stress (about 1.9 GPa), indicating that the temperature dependence of yield (flow) stress is not evident before glass transition upon heating. During the glass transition, the flow stress drops abruptly and decreases to be almost zero around the end temperature of glass transition (T g-end ). Around the glass transition region, the decrease in the flow temperature with an increasing applied stress (or the abrupt decrease in flow stress upon heating) can be roughly attributed to the drastic change in viscosity during glass transition. It should be pointed out that the experimental observation shown in this study might obscure an important phenomenon that the glass transition may be shifted under this circumstance. Since an abrupt (but continuous) change in specific heat and thermal expansion occurs at T g, commonly the glass transition is simply treated to be a second order phase transition, on which the pressure should have little effect since it does not involve discontinuous change in volume. However, sharp glass transition has never been observed on laboratory time scale, indicating that the experimentally observable glass transition is not a thermodynamic equilibrium phase transformation [16,17]. Via the enthalpy recovery method, Samwer et al. [18] derived an increase of glass transition temperature with hydrostatic pressure of about 3.6 K/GPa in a Zr 46:25 Ti 8:25 Cu 7:5 Ni 10:0 Be 27:5 amorphous alloy (Johnson glass Vit 4). Nevertheless, the effect of pressure on the glass transition temperature has not been observed directly in the metallic systems, and the effect of the stress and deformation on glass transition temperature is still unknown. It has been reported that the deformation at temperatures close to T g can induce structural disorder and free volume increase [19,20]. In terms of a model suggested by van den Beuckle et al., the glass transition is described to be the continuous approach of free volume towards equilibrium upon warming up [21]. Therefore, the increase of free volume induced by deformation may result from the modified equilibrium between production and annihilation rate of free volume, which would result in the shift of glass transition temperature that may make a contribution to the flow temperature reduction observed in this study. This cannot be analyzed further in this work because direct evidence is still lacking. 4. Conclusions In summary, for a bulk Zr 41:2 Ti 13:8 Cu 12:5 - Ni 10:0 Be 22:5 metallic glass, the flow temperature was reduced by 45 K/1.37 GPa under compression stress and the processible temperature range is expected to be possibly widened by about 70 K/1.37 GPa, while a heating rate of 5 K/min was applied.

5 H.J. Jin et al. / Scripta Materialia 47 (2002) The flow stress was found to decrease drastically during glass transition and to approach almost zero around the end temperature of glass transition. Acknowledgements The financial supports from the National Natural Science Foundation of China (grant no ) and the Max-Planck Society of Germany are acknowledged. References [1] Busch B, Bakke E, Johnson WL. Acta Mater 1998;46:4725. [2] Inoue A, Kawamura Y, Shibata T, Sasamori K. Mater Trans JIM 1996;37:1337. [3] Inoue A, Horikiri H, Kato A, Masumoto T. Mater Trans JIM 1994;35:79. [4] Zhou F, Zhang XH, Lu K. J Mater Res 1998;13:784. [5] Kawamura Y, Kato H, Inoue A. Appl Phys Lett 1995;67:2008. [6] Gu XJ, Jin HJ, Zhang HW, Wang JQ, Lu K. Scripta Mater 2001;45:1091. [7] Ye F, Lu K. Acta Mater 1999;47:2449. [8] Argon AS. Acta Metall Mater 1979;27:47. [9] Nieh TG, Wadsworth J, Liu CT, Ohkubo T, Hirotsu Y. Acta Mater 2001;49:2887. [10] Reger-Leonhard A, Heilmaier M, Eckert J. Scripta Mater 2000;43:459. [11] Myung W-N, Bae H-Y, Hwang I-S, Kim H-G, Nishiyama N, Inoue A, et al. Mater Sci Eng A 2001; :687. [12] Nishiyama N, Inoue A. Mater Trans JIM 1999;40:64. [13] Jiang JZ, Zhou TJ, Rasmussen H. Appl Phys Lett 2000;77:3553. [14] Burning R, Samwer K. Phys Rev B 1992;46: [15] Bruck HA, Christman T, Rosakis AJ, Johnson WL. Scripta Metal Mater 1994;30:429. [16] Kauzman W. Chem Rev 1948;43:219. [17] J ackle J. Rep Prog Phys 1986;49:171. [18] Samwer K, Busch R, Johnson WL. Phys Rev Lett 1999;82:580. [19] de Hey P, Sietsma J, van den Beukel A. Acta Mater 1998;46:5873. [20] Chen HS, Kato H, Inoue A, Saida J, Nishiyama N. Appl Phys Lett 2001;79:60. [21] van den Beukel A, Sietsma J. Acta Metall Mater 1990;38:383.