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Mechanical stability of Ca65Mg15Zn20 bulk metallic glass during deformation in the supercooled liquid region
Materials Science and Engineering A 480 (2008) 198–204
Mechanical stability of Ca65Mg15Zn20 bulk metallic glass during deformation in the supercooled liquid region K.J. Laws, B. Gun, M. Ferry ∗ ARC Centre of Excellence for Design in Light Metals & School of Materials Science and Engineering, University of New South Wales, Sydney 2052, Australia Received 22 February 2007; received in revised form 28 June 2007; accepted 12 July 2007
Abstract Ca65 Mg15 Zn20 bulk metallic glass (BMG) samples of dimensions 3.2 mm × 7 mm × 125 mm were prepared using a low-pressure die casting technique. These samples were ground to produce tensile test pieces in compliance with ASTM E8-04. This work is the first reported study of the tensile behaviour of Ca65 Mg15 Zn20 BMG in the supercooled liquid region (105–120 ◦ C). Two deformation conditions were used for testing: (i) constant strain rate testing from 10−3 to 10−4 s−1 and (ii) constant load testing using loads of 20–50 N applied to a tensile sample during heating at a constant rate of 5 ◦ C s−1 . The maximum elongation to failure in the BMG was in excess of 850% for constant load testing although, under isothermal testing conditions, most samples failed after ∼200% elongation. It is concluded that large superplastic elongations (>500%) during isothermal tensile straining is difficult in this alloy due to the onset of crystallization. © 2007 Elsevier B.V. All rights reserved. Keywords: Bulk metallic glass; Mechanical behaviour; Tensile testing; Crystallization
1. Introduction Calcium-based bulk metallic glasses (BMGs) are a relatively new and unique group of lightweight amorphous alloys. The first Ca-based BMGs were reported by Amiya and Inoue in 2002, where they produced, by conventional copper mould casting, fully amorphous rods with a maximum diameter of 4 mm for two ternary alloys, Ca57 Mg19 Cu24 and Ca60 Mg20 Ag20 [1], and 7 mm for a quaternary alloy, Ca60 Mg20 Ag10 Cu10 [2]. More recently, numerous multi-component glass forming systems based on Ca have been explored, such as Ca–Mg–Zn, Ca–Mg–Al, Ca–Al–Cu and various combinations of these systems [3–7]. Ca-based BMGs are potential candidates for a range of applications as they exhibit a high glass forming ability/stability and unique physical and chemical properties. For example, most Ca-based crystalline alloys oxidize in air in a matter of days, but many Ca-based metallic glasses maintain a metallic lustre for long periods after casting [8]. Other useful properties include a low density (∼2000 kg/m3 ), relatively high compressive yield strengths (∼350 MPa), low
Young’s modulus (∼17–20 GPa), low glass transition temperature (Tg ∼ 110 ◦ C), and a large supercooled liquid (SCL) region (TX ∼ 30–70 ◦ C) where superplastic flow is possible. A dramatic increase in plasticity at temperatures above the glass transition temperature (Tg ) is characteristic of many amorphous alloys and is usually explained by the transformation of the glassy state into a supercooled liquid (SCL) state, where a dramatic decrease in viscosity is observed [9]. Studies of the flow behaviour of BMGs in the SCL region have mainly been carried out by compressive testing with only a limited number of studies reported on tensile deformation (see e.g. Ref. [10]). There are no reports of the tensile flow behaviour of Ca-based BMGs. In the present work, standardised tensile test specimens, in compliance with ASTM E8-04, have been produced to investigate the flow behaviour of Ca65 Mg15 Zn20 BMG in the SCL region. 2. Experimental procedure 2.1. Casting and preparation of tensile samples
∗
Corresponding author. Tel.: +61 2 9385 4453. E-mail address: [email protected] (M. Ferry).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.07.034
Ca65 Mg15 Zn20 bulk metallic glass samples of dimensions 3.2 mm × 7 mm × 125 mm were prepared using a repetitive
K.J. Laws et al. / Materials Science and Engineering A 480 (2008) 198–204
low-pressure die casting technique [11]. Using this caster, a two-stage melting and die casting procedure was carried out for generating the samples. A mixture of pure elements and master alloys, namely Ca (99.8 wt.%), 80 wt.% Ca/20 wt.% Mg master alloy (99.8 wt.% purity) and Zn (99.9 wt.%) were used for charge preparation. All materials were mechanically scrubbed immediately prior to melting. The balance of Ca was placed in the bottom of a Mo crucible, followed by the balance of Zn and approximately one-third the balance of master alloy. The melting chamber was vacuum purged with argon gas (99.997 vol.%) and left for 20 min under a circulating argon atmosphere. The charge was induction heated to 700 ◦ C, held for 5 min, stirred with a tungsten rod then cooled to room temperature. The remainder of the master alloy was added to the crucible, reheated to 700 ◦ C, held for 5 min and stirred. The charge was cooled to the desired injection temperature, where it was held for a further 2.5 min and stirred immediately prior to injection casting. Further details of this casting facility are given elsewhere [11]. The composition of the as-cast samples was analysed by both electron probe microanalysis (EPMA) and X-ray fluorescence (XRF) and found to be Ca65 Mg15 Zn20 ± 0.28 at.%. Due to the brittle nature of the alloy at room temperature [8], high quality tensile test samples, in accordance with ASTM E8-04, were produced from the injection-cast rods using an inhouse grinding technique involving the simultaneous rotation and grinding of the rectangular-shaped sample to generate 3 mm diameter samples of gauge length 12.3 mm.
199
Fig. 1. Schematic representation of the apparatus used for constant strain and constant load tensile testing of Ca65 Mg15 Zn20 BMG samples.
2.2. Mechanical testing Due to the reactive nature of the alloy at elevated temperatures, the tensile samples were deformed at a given temperature in silicone oil using an MTS 810 servo-hydraulic testing machine equipped with a 1 kN load cell. Fig. 1 is a schematic representation of the experimental apparatus used for both constant
Fig. 2. (a) Complete DSC profile of the as-cast BMG at a constant heating rate of 5 ◦ C/min. Magnification of: (b) Tg region, (c) TX region and (d) Tm and TL regions.
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strain and constant load testing. During testing, a type K thermocouple was positioned in close proximity to the specimen to monitor the temperature at the specimen surface which maintained constant temperature within ±1 ◦ C. The chamber was preheated prior to mounting the specimens to minimise heat up time. Due to the fact that Tg is not a unique physical property of the glass but rather a rate dependent kinetic quantity, the heating rate of the specimen was kept constant at ∼5 ◦ C/min after reaching 90 ◦ C. Tensile testing commenced only after the sample test temperature was stabilized for ∼2 min. Two types of tensile testing experiments were carried out: (i) constant strain rate testing (10−3 –10−4 s−1 ) at constant temperature in the range 105–120 ◦ C and (ii) constant load testing (20–50 N) during heating at a rate of 5 ◦ C/min. In this work, over 50 tensile test samples were used with each test condition repeated up to four times for consistency. The degree of crystallization of both the as-cast and deformed samples was determined by X-ray diffraction (XRD) using a Philips MRD diffractometer using Cu K␣ radiation. The thermal properties of the alloy were determined using a TA 2010 differential scanning calorimeter (DSC) at a heating rate of 5 ◦ C/min. For this heating rate, the error in determining Tg and the crystallization onset temperature (TX ) was less than 0.5 ◦ C using indium as a calibration standard. Isothermal DSC measurements have recently been carried out on the same material at various test temperatures to determine the onset of crystallization and the kinetics of crystallization.
Fig. 3. Isothermal DSC profiles for the as-cast BMG for temperatures in the range: (a) 135–125 ◦ C and (b) 120–110 ◦ C (temperatures were approached at a heating rate of 5 ◦ C/min).
3. Results and discussion 3.1. DSC evaluation The thermal properties of the BMG, such as Tg , TX , melting temperature (Tm ) and liquidus temperature (TL ), were determined on 3 mm sections of the as-cast material using DSC at a heating rate of 5 ◦ C/min: this also replicates the heating rate used to approach each testing temperature in the constant strain rate tests as well as the constant load tests. For the present alloy: Tg = 105 ◦ C, TX = 140 ◦ C, T (SCL) = 35 ◦ C, Tm = 357 ◦ C and TL = 337 ◦ C (Fig. 2). Isothermal DSC was also carried out at temperatures ranging from 110 to 135 ◦ C for determining the allowable time frame for tensile testing before the onset of crystallization. Isothermal DSC traces are given in Fig. 3 for the temperature range 110–135 ◦ C [12]. It can be seen that holding at a temperature of 125 ◦ C or greater (greater than half the SCL region) results in a very short incubation period prior to the onset of crystallization. This suggests that only half of the SCL region is of practical use in low strainrate laboratory testing or manufacture of components if a fully amorphous structure is to be maintained. This behaviour is expected since the size of the SCL region TX ∼ 35 ◦ C (which is quite small relative to other BMGs with such a high glass forming ability [13,14]) is a good indicator of glass stability [14,15]. Using isothermal DSC data, a simple transformation diagram was constructed with respect to the onset and completion times
of crystallization (Fig. 4). The boundaries between these regions are given as t0.05 and t0.95 , as determined from Fig. 3. The diagram shows the common trend, whereby the allowable time for retaining the amorphous phase decreases with increasing temperature: this information is useful for determining suitable test times for tensile testing with respect to the onset of crystallization.
Fig. 4. Transformation diagram of the BMG generated from isothermal DSC data showing the amorphous, semi-crystalline and crystalline phase fields.
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Table 1 Summary of tensile test results including either maximum flow stress (σ f ) (brittle failure) or peak stress (σ p ) (ductile failure), maximum strain % (εmax ), test time to failure (tf ) and the time delay of the stress response due to crystallization (tX )
3.2. Constant strain rate tensile testing The isothermal DSC test results confirmed that the Ca65 Mg15 Zn20 alloy was relatively unstable and crystallized rapidly even at relatively low temperatures within the SCL region (Figs. 3 and 4). Consequently, tensile testing was carried out at temperatures in the range 105–120 ◦ C at increments of 5 ◦ C. The experiments indicated that this alloy was highly strain rate sensitive and that the maximum applied strain rate achievable in tension was only 1 × 10−3 s−1 (which is quite low when compared to other BMG systems [16–19]). Hence, tensile testing was carried out at strain rates ranging from 1 × 10−3 to 1 × 10−4 s−1 . Table 1 gives various test results including flow stress (σ f ) or peak stress (σ p ), maximum percentage strain (εmax ), time to failure after reaching the test temperature (tf ) and the time delay of the stress response due to crystallization (tX ) compared to an unstrained sample. The shaded values in Table 1 indicate samples that have failed in a brittle manner. With the exception of a single sample, all samples failed in either a brittle manner (εmax < 2.3%) or continued to flow plastically until failure. The failure mechanism of the sample tested at 105 ◦ C and 10−4 s−1 may be an extreme example of non-Newtonian flow, which would be expected since testing was carried out so close to Tg . For the samples that deformed by homogeneous plastic flow, a typical stress overshoot and yield drop to a much lower stress was observed in most cases; this behaviour is similar to that reported for other alloy systems [16–20]. This strain-induced softening after an initial peak in the stress is generally explained by an increase in the free volume of the amorphous structure and the generation of shear bands, which may act as dislocation-like defects, both of which contribute to a lower flow stress [9,21]. From Table 1, there is a clear increase in peak stress (σ p ) with
decreasing test temperature (for a given strain rate), which is analogous to an increase in viscosity of the supercooled liquid. The flow behaviour in the SCL region was examined with respect to flow stress and strain rate with results given in Fig. 5. For the highest test temperature, there is a reasonably linear relationship between log stress and log strain rate; such behaviour is often described by the relation [22]: σ = k˙εm
(1)
where m is the strain-rate sensitivity (m-value). It can be seen in Fig. 5 that m-values are close to unity for much of the dataset,
Fig. 5. Stress as a function of strain rate (log–log format) during isothermal tensile testing in the range 105–120 ◦ C. The slope of each line indicates the strain-rate sensitivity of the material (m).
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Fig. 6. The effect of time on tensile stress during holding samples at: (a) 120 ◦ C and strained at 1 × 10−3 s−1 and 7.5 × 10−4 s−1 , (b) 120 ◦ C and strained at (5–1) × 10−4 s−1 , (c) 115 ◦ C and strained at (5–1) × 10−4 s−1 and (d) 110 ◦ C and strained at (5–1) × 10−4 s−1 .
which indicates that plastic flow occurs largely in a Newtonian manner [22]. However, there is a clear transition from Newtonian to non-Newtonian flow, which corresponds to a gradual decrease in the m-value at lower temperatures in the SCL region. Two theories have been proposed for explaining the transition from Newtonian to non-Newtonian behaviour [20]: (i) transition state theory of intrinsic plastic flow of glass [23] and (ii) concurrent nanocrystallization [24]. With respect to the latter, Nieh et al. [24] proposed that non-Newtonian behaviour is associated with stress-driven nanocrystallization during deformation. This suggests that when a metallic glass contains aggregates of nanocrystals dispersed in an amorphous structure, it deforms by grain boundary sliding, where m-values of less than unity are expected [20]. The maximum ductility in the present BMG under isothermal straining conditions was ∼600% elongation which is reasonable for a metallic glass. It is also important to note that the samples tested at 120 ◦ C and strain rates of 1 × 10−3 s−1 and 7.5 × 10−4 s−1 exhibited poor ductility (less elongation than expected) due to localized necking. With respect to the maximum obtainable strain, the role of crystallization is quite apparent. With the exception of the tests carried out at 120 ◦ C
and strain rates of 1 × 10−3 s−1 and 7.5 × 10−4 s−1 (where the time for crystallization is not reached before failure (Fig. 6a)), the maximum strain achieved by a sample is proportional to the duration of the test. Hence, sample elongations tend to halve as strain rate is halved (Table 1). It has been documented for various BMG systems [16–20] that an increase in stress is associated with temperature–timeinduced crystallization; this is also likely to be the case for the present alloy. These results also indicate that the strength of the partially crystalline phases is higher than that of a completely amorphous phase at elevated temperatures. The delay in the onset time of the stress increase due to crystallization (tX ) can be seen in Fig. 6b–d, showing that tX varies with both temperature and strain rate. Included in Fig. 6 are DSC profiles representing the onset of crystallization of an unstrained sample. Here, higher temperatures yield smaller tX values and higher strain rates yield larger tX values. Similar behaviour has been documented elsewhere [17,18,20]. This time discrepancy in the onset of the stress increase due to crystallization suggests that, under appropriate deformation conditions, a degradation of the crystallization reaction occurs, i.e. deformation-induced stabilization of the supercooled liquid.
K.J. Laws et al. / Materials Science and Engineering A 480 (2008) 198–204 Table 2 Summary of constant load tensile test results including maximum percentage strain (εmax ), maximum strain rate (˙εmax ), test time to failure (tf ) and the temperature at failure (Tf ) Load (N) 50 40 20
max (%)
ε˙ max (s−1 )
tf (min)
Tf (◦ C)
853 517 260
0.31 0.27 0.10
10.7 11.9 12.6
144 145 147
Fig. 7. Sample tensile deformed under a constant load of 50 N at a heating rate of 5 ◦ C/min. The elongation at failure was 850%.
The mechanism for this deformation-induced stabilization is not clear, although it is suggested that it may be due to the suppression of an incipient-stage reaction for crystallization that, at higher strain rates, is a result of a rapid increase in the flow volume generated during deformation [15–18]. 3.3. Constant load tensile testing In a series of experiments, samples were deformed under a constant load ranging from 20 to 50 N during heating at a rate of 5 ◦ C/min for investigating the temperature-dependence of plastic flow and the variation in strain rate obtained above the glass transition temperature. The results are given in Table 2. The variation in strain rate was taken with respect to the original gauge length with strain given as percentage elongation. During testing under constant load, strain rates greater than 1 × 10−1 s−1 and elongations to failure in excess of 850% were achieved (Fig. 7). A plot of instantaneous strain rate under constant load is given in Fig. 8 which includes DSC data of an unstrained sample for the same temperature interval and heating rate (5 ◦ C/min). The data show extremely low strain rates at temperatures below Tg , followed by a steady increase through the SCL region up to TX . Beyond TX there is a more rapid increase in strain rate and eventual failure.
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In the very advanced stages of straining (i.e. high deformation temperatures), a part of the increase in strain rate may be attributed to localized thinning of the sample that, under constant load, results in an increase in stress. However, Fig. 8 shows a near linear trend between strain rate, applied load and temperature throughout the SCL region. It is reasonable to assume that this linear trend shown in Fig. 8 is caused by the change in viscosity of the BMG as temperature is increased. The equilibrium viscosity of a BMG is often explained using the Vogel–Fulcher–Tammann (VFT) relation. In the Newtonian regime (low strain rates), this relation is given in the form: DT0 η = η0 exp (2) T − T0 where T0 is the VFT temperature, D termed the fragility parameter and η0 is the high temperature limit of viscosity. It is also known that viscosity in the Newtonian regime may be described by flow stress (σ f ) and strain rate [23] such that: η=
σf 3˙ε
Combining Eqs. (2) and (3) and rearranging gives: DT0 σf − ln ε˙ = ln 3η0 T − T0
(3)
(4)
A plot of ln(˙ε) as a function 1/T − T0 is expected to yield straight, parallel lines of gradient DT0 and an intercept related to σ f : this explains the results given in Fig. 8. As may be expected, higher strains in the SCL region results in a deviation from linearity due to localized thinning of the samples, although it may also be a cumulative effect of the viscosity decreasing with the increasing strain rate as flow conditions change from Newtonian to non-Newtonian due to an increase in sample free volume. As the temperature is increased above TX , a marked deviation in the flow rate between samples becomes apparent (Fig. 8). For the sample loaded at 20 N, there is a marked decrease in the rate at which the strain rate is increasing compared to those samples under higher load. Such deviation in flow behaviour is shown by an inflexion in the temperature/strain-rate curve at ∼143.4 ◦ C. Although not as pronounced, the sample loaded at 40 N also shows similar behaviour with an inflexion at ∼144.5 ◦ C immediately before failure. In comparison to the case of the sample held at 50 N, this information further supports the theory of deformation-induced stabilization of the supercooled liquid, whereby the strengthening response of crystallization has been delayed or completely suppressed for samples under higher loads and higher rate of deformation. Here it can be appreciated that this deformation method shows potential of greater elongations before failure due to crystallization hardening. 4. Conclusions
Fig. 8. Diagram showing instantaneous strain rates as a function of temperature during constant load testing. Included in the diagram is the DSC profile of an unstrained sample for the same heating rate and temperature interval.
The tensile flow behaviour of a Ca65 Mg15 Zn20 bulk metallic glass (BMG) was investigated in the supercooled liquid (SCL) region for a range of strain rates (10−3 –10−4 s−1 ) and temperatures (105–120 ◦ C). It was found that:
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• the flow stress increases to a maximum followed by a decrease in stress to a low steady state value; • a maximum elongation of ∼600% is achieved when testing at 110 ◦ C for a strain rate of 5 × 10−4 s−1 ; • due to the onset of crystallization in the SCL region, there is an increase in stress which indicates that, at elevated temperatures, the overall strength of the semi-crystalline material is higher than that of the fully amorphous material and • the stress increase associated with crystallization varies with test temperature and strain rate, such that the undesirable crystallization reaction is more readily suppressed at lower temperatures and higher strain rates. The tensile flow behaviour of the BMG was also investigated under conditions of constant load (20–50 N) and constant heating rate (5 ◦ C/min). It was found that: • an increase in the deformation temperature in the SCL region results in a concomitant increase in strain rate of the material which is be explained by current deformation and viscosity theory (i.e. Eqs. (2) and (3)) and • hardening associated with the onset of crystallization is more readily suppressed at higher loads/deformation rates which allows higher tensile strains (>800%) to be achieved. Acknowledgement The authors would like to thank the Australian Research Council (ARC) for partial funding of this work via the ARC Centre of Excellence for Design in Light Metals (CEO561574).
References [1] K. Amiya, A. Inoue. Mater. Trans. JIM 43 (2002) 81. [2] K. Amiya, A. Inoue. Mater. Trans. JIM 43 (2002) 2578. [3] O.N. Senkov, J.M. Scott, MRS Proceedings, vol. 806, Materials Research Society, Warrendale, PA, 2003, p. 145. [4] O.N. Senkov, J.M. Scott, Mater. Lett. 58 (2004) 1375. [5] O.N. Senkov, J.M. Scott, Scripta Mater. 50 (2004) 449. [6] F.Q. Guo, S.J. Poon, G.J. Shiflet, Appl. Phys. Lett. 84 (2003) 37. [7] E.S. Park, D.H. Kim, J. Mater. Res. 19 (2004) 685. [8] O.N. Senkov, D.B. Miracle, J.M. Scott, Intermetallics 14 (2006) 1055. [9] F. Spaepen, Acta Metall. 25 (1977) 407. [10] G. Wang, J. Shen, J.F. Sun, Y.J. Huang, J. Zou, Z.P. Lu, Z.H. Stachurski, B.D. Zhou, J. Non-Cryst. Solids 351 (2005) 209. [11] K.J. Laws, B. Gun, M. Ferry, Mater. Sci. Eng. A 475 (2008) 348. [12] K.J. Laws, PhD Thesis, University of New South Wales (2007). [13] G.J. Fan, H. Choo, P.K. Liaw, J. Non-Cryst. Solids 353 (2007) 102. [14] W.Y. Liu, H.F. Zhang, A.M. Wang, H. Li, Z.Q. Hu, Mater. Sci. Eng. A 459 (2007) 196. [15] A. Inoue, Acta Mater. 48 (2000) 279. [16] B. Gun, K.J. Laws, M. Ferry, J. Non-Cryst. Solids 352 (2006) 3896. [17] A. Inoue, Y. Kawamura, Y. Saotome, Mater. Sci. Forum 147 (1997) 233. [18] Y. Kawamura, T. Shibata, A. Inoue, Mater. Trans. JIM 40 (1999) 335. [19] B. Gun, K.J. Laws, M. Ferry, J. Non-Cryst. Solids 352 (2006) 3887. [20] W.J. Kim, Intermetallics 15 (2007) 282. [21] T. Masumoto, R. Maddin, Mater. Sci. Eng. 19 (1975) 1. [22] A. Inoue, Y. Saotome, Metals 3 (1993) 51. [23] M. Bletry, P. Guyot, Y. Brechet, J.J. Blandin, J.L. Soubeyroux, Mater. Sci. Eng. A 387–389 (2004) 1005. [24] T.G. Nieh, T. Mukai, C.T. Liu, J. Wadsworth, Scripta Mater. 40 (1999) 1021.
Mechanical stability of Ca65Mg15Zn20 bulk metallic glass during deformation in the supercooled liquid region K.J. Laws, B. Gun, M. Ferry ∗ ARC Centre of Excellence for Design in Light Metals & School of Materials Science and Engineering, University of New South Wales, Sydney 2052, Australia Received 22 February 2007; received in revised form 28 June 2007; accepted 12 July 2007
Abstract Ca65 Mg15 Zn20 bulk metallic glass (BMG) samples of dimensions 3.2 mm × 7 mm × 125 mm were prepared using a low-pressure die casting technique. These samples were ground to produce tensile test pieces in compliance with ASTM E8-04. This work is the first reported study of the tensile behaviour of Ca65 Mg15 Zn20 BMG in the supercooled liquid region (105–120 ◦ C). Two deformation conditions were used for testing: (i) constant strain rate testing from 10−3 to 10−4 s−1 and (ii) constant load testing using loads of 20–50 N applied to a tensile sample during heating at a constant rate of 5 ◦ C s−1 . The maximum elongation to failure in the BMG was in excess of 850% for constant load testing although, under isothermal testing conditions, most samples failed after ∼200% elongation. It is concluded that large superplastic elongations (>500%) during isothermal tensile straining is difficult in this alloy due to the onset of crystallization. © 2007 Elsevier B.V. All rights reserved. Keywords: Bulk metallic glass; Mechanical behaviour; Tensile testing; Crystallization
1. Introduction Calcium-based bulk metallic glasses (BMGs) are a relatively new and unique group of lightweight amorphous alloys. The first Ca-based BMGs were reported by Amiya and Inoue in 2002, where they produced, by conventional copper mould casting, fully amorphous rods with a maximum diameter of 4 mm for two ternary alloys, Ca57 Mg19 Cu24 and Ca60 Mg20 Ag20 [1], and 7 mm for a quaternary alloy, Ca60 Mg20 Ag10 Cu10 [2]. More recently, numerous multi-component glass forming systems based on Ca have been explored, such as Ca–Mg–Zn, Ca–Mg–Al, Ca–Al–Cu and various combinations of these systems [3–7]. Ca-based BMGs are potential candidates for a range of applications as they exhibit a high glass forming ability/stability and unique physical and chemical properties. For example, most Ca-based crystalline alloys oxidize in air in a matter of days, but many Ca-based metallic glasses maintain a metallic lustre for long periods after casting [8]. Other useful properties include a low density (∼2000 kg/m3 ), relatively high compressive yield strengths (∼350 MPa), low
Young’s modulus (∼17–20 GPa), low glass transition temperature (Tg ∼ 110 ◦ C), and a large supercooled liquid (SCL) region (TX ∼ 30–70 ◦ C) where superplastic flow is possible. A dramatic increase in plasticity at temperatures above the glass transition temperature (Tg ) is characteristic of many amorphous alloys and is usually explained by the transformation of the glassy state into a supercooled liquid (SCL) state, where a dramatic decrease in viscosity is observed [9]. Studies of the flow behaviour of BMGs in the SCL region have mainly been carried out by compressive testing with only a limited number of studies reported on tensile deformation (see e.g. Ref. [10]). There are no reports of the tensile flow behaviour of Ca-based BMGs. In the present work, standardised tensile test specimens, in compliance with ASTM E8-04, have been produced to investigate the flow behaviour of Ca65 Mg15 Zn20 BMG in the SCL region. 2. Experimental procedure 2.1. Casting and preparation of tensile samples
∗
Corresponding author. Tel.: +61 2 9385 4453. E-mail address: [email protected] (M. Ferry).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.07.034
Ca65 Mg15 Zn20 bulk metallic glass samples of dimensions 3.2 mm × 7 mm × 125 mm were prepared using a repetitive
K.J. Laws et al. / Materials Science and Engineering A 480 (2008) 198–204
low-pressure die casting technique [11]. Using this caster, a two-stage melting and die casting procedure was carried out for generating the samples. A mixture of pure elements and master alloys, namely Ca (99.8 wt.%), 80 wt.% Ca/20 wt.% Mg master alloy (99.8 wt.% purity) and Zn (99.9 wt.%) were used for charge preparation. All materials were mechanically scrubbed immediately prior to melting. The balance of Ca was placed in the bottom of a Mo crucible, followed by the balance of Zn and approximately one-third the balance of master alloy. The melting chamber was vacuum purged with argon gas (99.997 vol.%) and left for 20 min under a circulating argon atmosphere. The charge was induction heated to 700 ◦ C, held for 5 min, stirred with a tungsten rod then cooled to room temperature. The remainder of the master alloy was added to the crucible, reheated to 700 ◦ C, held for 5 min and stirred. The charge was cooled to the desired injection temperature, where it was held for a further 2.5 min and stirred immediately prior to injection casting. Further details of this casting facility are given elsewhere [11]. The composition of the as-cast samples was analysed by both electron probe microanalysis (EPMA) and X-ray fluorescence (XRF) and found to be Ca65 Mg15 Zn20 ± 0.28 at.%. Due to the brittle nature of the alloy at room temperature [8], high quality tensile test samples, in accordance with ASTM E8-04, were produced from the injection-cast rods using an inhouse grinding technique involving the simultaneous rotation and grinding of the rectangular-shaped sample to generate 3 mm diameter samples of gauge length 12.3 mm.
199
Fig. 1. Schematic representation of the apparatus used for constant strain and constant load tensile testing of Ca65 Mg15 Zn20 BMG samples.
2.2. Mechanical testing Due to the reactive nature of the alloy at elevated temperatures, the tensile samples were deformed at a given temperature in silicone oil using an MTS 810 servo-hydraulic testing machine equipped with a 1 kN load cell. Fig. 1 is a schematic representation of the experimental apparatus used for both constant
Fig. 2. (a) Complete DSC profile of the as-cast BMG at a constant heating rate of 5 ◦ C/min. Magnification of: (b) Tg region, (c) TX region and (d) Tm and TL regions.
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strain and constant load testing. During testing, a type K thermocouple was positioned in close proximity to the specimen to monitor the temperature at the specimen surface which maintained constant temperature within ±1 ◦ C. The chamber was preheated prior to mounting the specimens to minimise heat up time. Due to the fact that Tg is not a unique physical property of the glass but rather a rate dependent kinetic quantity, the heating rate of the specimen was kept constant at ∼5 ◦ C/min after reaching 90 ◦ C. Tensile testing commenced only after the sample test temperature was stabilized for ∼2 min. Two types of tensile testing experiments were carried out: (i) constant strain rate testing (10−3 –10−4 s−1 ) at constant temperature in the range 105–120 ◦ C and (ii) constant load testing (20–50 N) during heating at a rate of 5 ◦ C/min. In this work, over 50 tensile test samples were used with each test condition repeated up to four times for consistency. The degree of crystallization of both the as-cast and deformed samples was determined by X-ray diffraction (XRD) using a Philips MRD diffractometer using Cu K␣ radiation. The thermal properties of the alloy were determined using a TA 2010 differential scanning calorimeter (DSC) at a heating rate of 5 ◦ C/min. For this heating rate, the error in determining Tg and the crystallization onset temperature (TX ) was less than 0.5 ◦ C using indium as a calibration standard. Isothermal DSC measurements have recently been carried out on the same material at various test temperatures to determine the onset of crystallization and the kinetics of crystallization.
Fig. 3. Isothermal DSC profiles for the as-cast BMG for temperatures in the range: (a) 135–125 ◦ C and (b) 120–110 ◦ C (temperatures were approached at a heating rate of 5 ◦ C/min).
3. Results and discussion 3.1. DSC evaluation The thermal properties of the BMG, such as Tg , TX , melting temperature (Tm ) and liquidus temperature (TL ), were determined on 3 mm sections of the as-cast material using DSC at a heating rate of 5 ◦ C/min: this also replicates the heating rate used to approach each testing temperature in the constant strain rate tests as well as the constant load tests. For the present alloy: Tg = 105 ◦ C, TX = 140 ◦ C, T (SCL) = 35 ◦ C, Tm = 357 ◦ C and TL = 337 ◦ C (Fig. 2). Isothermal DSC was also carried out at temperatures ranging from 110 to 135 ◦ C for determining the allowable time frame for tensile testing before the onset of crystallization. Isothermal DSC traces are given in Fig. 3 for the temperature range 110–135 ◦ C [12]. It can be seen that holding at a temperature of 125 ◦ C or greater (greater than half the SCL region) results in a very short incubation period prior to the onset of crystallization. This suggests that only half of the SCL region is of practical use in low strainrate laboratory testing or manufacture of components if a fully amorphous structure is to be maintained. This behaviour is expected since the size of the SCL region TX ∼ 35 ◦ C (which is quite small relative to other BMGs with such a high glass forming ability [13,14]) is a good indicator of glass stability [14,15]. Using isothermal DSC data, a simple transformation diagram was constructed with respect to the onset and completion times
of crystallization (Fig. 4). The boundaries between these regions are given as t0.05 and t0.95 , as determined from Fig. 3. The diagram shows the common trend, whereby the allowable time for retaining the amorphous phase decreases with increasing temperature: this information is useful for determining suitable test times for tensile testing with respect to the onset of crystallization.
Fig. 4. Transformation diagram of the BMG generated from isothermal DSC data showing the amorphous, semi-crystalline and crystalline phase fields.
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Table 1 Summary of tensile test results including either maximum flow stress (σ f ) (brittle failure) or peak stress (σ p ) (ductile failure), maximum strain % (εmax ), test time to failure (tf ) and the time delay of the stress response due to crystallization (tX )
3.2. Constant strain rate tensile testing The isothermal DSC test results confirmed that the Ca65 Mg15 Zn20 alloy was relatively unstable and crystallized rapidly even at relatively low temperatures within the SCL region (Figs. 3 and 4). Consequently, tensile testing was carried out at temperatures in the range 105–120 ◦ C at increments of 5 ◦ C. The experiments indicated that this alloy was highly strain rate sensitive and that the maximum applied strain rate achievable in tension was only 1 × 10−3 s−1 (which is quite low when compared to other BMG systems [16–19]). Hence, tensile testing was carried out at strain rates ranging from 1 × 10−3 to 1 × 10−4 s−1 . Table 1 gives various test results including flow stress (σ f ) or peak stress (σ p ), maximum percentage strain (εmax ), time to failure after reaching the test temperature (tf ) and the time delay of the stress response due to crystallization (tX ) compared to an unstrained sample. The shaded values in Table 1 indicate samples that have failed in a brittle manner. With the exception of a single sample, all samples failed in either a brittle manner (εmax < 2.3%) or continued to flow plastically until failure. The failure mechanism of the sample tested at 105 ◦ C and 10−4 s−1 may be an extreme example of non-Newtonian flow, which would be expected since testing was carried out so close to Tg . For the samples that deformed by homogeneous plastic flow, a typical stress overshoot and yield drop to a much lower stress was observed in most cases; this behaviour is similar to that reported for other alloy systems [16–20]. This strain-induced softening after an initial peak in the stress is generally explained by an increase in the free volume of the amorphous structure and the generation of shear bands, which may act as dislocation-like defects, both of which contribute to a lower flow stress [9,21]. From Table 1, there is a clear increase in peak stress (σ p ) with
decreasing test temperature (for a given strain rate), which is analogous to an increase in viscosity of the supercooled liquid. The flow behaviour in the SCL region was examined with respect to flow stress and strain rate with results given in Fig. 5. For the highest test temperature, there is a reasonably linear relationship between log stress and log strain rate; such behaviour is often described by the relation [22]: σ = k˙εm
(1)
where m is the strain-rate sensitivity (m-value). It can be seen in Fig. 5 that m-values are close to unity for much of the dataset,
Fig. 5. Stress as a function of strain rate (log–log format) during isothermal tensile testing in the range 105–120 ◦ C. The slope of each line indicates the strain-rate sensitivity of the material (m).
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Fig. 6. The effect of time on tensile stress during holding samples at: (a) 120 ◦ C and strained at 1 × 10−3 s−1 and 7.5 × 10−4 s−1 , (b) 120 ◦ C and strained at (5–1) × 10−4 s−1 , (c) 115 ◦ C and strained at (5–1) × 10−4 s−1 and (d) 110 ◦ C and strained at (5–1) × 10−4 s−1 .
which indicates that plastic flow occurs largely in a Newtonian manner [22]. However, there is a clear transition from Newtonian to non-Newtonian flow, which corresponds to a gradual decrease in the m-value at lower temperatures in the SCL region. Two theories have been proposed for explaining the transition from Newtonian to non-Newtonian behaviour [20]: (i) transition state theory of intrinsic plastic flow of glass [23] and (ii) concurrent nanocrystallization [24]. With respect to the latter, Nieh et al. [24] proposed that non-Newtonian behaviour is associated with stress-driven nanocrystallization during deformation. This suggests that when a metallic glass contains aggregates of nanocrystals dispersed in an amorphous structure, it deforms by grain boundary sliding, where m-values of less than unity are expected [20]. The maximum ductility in the present BMG under isothermal straining conditions was ∼600% elongation which is reasonable for a metallic glass. It is also important to note that the samples tested at 120 ◦ C and strain rates of 1 × 10−3 s−1 and 7.5 × 10−4 s−1 exhibited poor ductility (less elongation than expected) due to localized necking. With respect to the maximum obtainable strain, the role of crystallization is quite apparent. With the exception of the tests carried out at 120 ◦ C
and strain rates of 1 × 10−3 s−1 and 7.5 × 10−4 s−1 (where the time for crystallization is not reached before failure (Fig. 6a)), the maximum strain achieved by a sample is proportional to the duration of the test. Hence, sample elongations tend to halve as strain rate is halved (Table 1). It has been documented for various BMG systems [16–20] that an increase in stress is associated with temperature–timeinduced crystallization; this is also likely to be the case for the present alloy. These results also indicate that the strength of the partially crystalline phases is higher than that of a completely amorphous phase at elevated temperatures. The delay in the onset time of the stress increase due to crystallization (tX ) can be seen in Fig. 6b–d, showing that tX varies with both temperature and strain rate. Included in Fig. 6 are DSC profiles representing the onset of crystallization of an unstrained sample. Here, higher temperatures yield smaller tX values and higher strain rates yield larger tX values. Similar behaviour has been documented elsewhere [17,18,20]. This time discrepancy in the onset of the stress increase due to crystallization suggests that, under appropriate deformation conditions, a degradation of the crystallization reaction occurs, i.e. deformation-induced stabilization of the supercooled liquid.
K.J. Laws et al. / Materials Science and Engineering A 480 (2008) 198–204 Table 2 Summary of constant load tensile test results including maximum percentage strain (εmax ), maximum strain rate (˙εmax ), test time to failure (tf ) and the temperature at failure (Tf ) Load (N) 50 40 20
max (%)
ε˙ max (s−1 )
tf (min)
Tf (◦ C)
853 517 260
0.31 0.27 0.10
10.7 11.9 12.6
144 145 147
Fig. 7. Sample tensile deformed under a constant load of 50 N at a heating rate of 5 ◦ C/min. The elongation at failure was 850%.
The mechanism for this deformation-induced stabilization is not clear, although it is suggested that it may be due to the suppression of an incipient-stage reaction for crystallization that, at higher strain rates, is a result of a rapid increase in the flow volume generated during deformation [15–18]. 3.3. Constant load tensile testing In a series of experiments, samples were deformed under a constant load ranging from 20 to 50 N during heating at a rate of 5 ◦ C/min for investigating the temperature-dependence of plastic flow and the variation in strain rate obtained above the glass transition temperature. The results are given in Table 2. The variation in strain rate was taken with respect to the original gauge length with strain given as percentage elongation. During testing under constant load, strain rates greater than 1 × 10−1 s−1 and elongations to failure in excess of 850% were achieved (Fig. 7). A plot of instantaneous strain rate under constant load is given in Fig. 8 which includes DSC data of an unstrained sample for the same temperature interval and heating rate (5 ◦ C/min). The data show extremely low strain rates at temperatures below Tg , followed by a steady increase through the SCL region up to TX . Beyond TX there is a more rapid increase in strain rate and eventual failure.
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In the very advanced stages of straining (i.e. high deformation temperatures), a part of the increase in strain rate may be attributed to localized thinning of the sample that, under constant load, results in an increase in stress. However, Fig. 8 shows a near linear trend between strain rate, applied load and temperature throughout the SCL region. It is reasonable to assume that this linear trend shown in Fig. 8 is caused by the change in viscosity of the BMG as temperature is increased. The equilibrium viscosity of a BMG is often explained using the Vogel–Fulcher–Tammann (VFT) relation. In the Newtonian regime (low strain rates), this relation is given in the form: DT0 η = η0 exp (2) T − T0 where T0 is the VFT temperature, D termed the fragility parameter and η0 is the high temperature limit of viscosity. It is also known that viscosity in the Newtonian regime may be described by flow stress (σ f ) and strain rate [23] such that: η=
σf 3˙ε
Combining Eqs. (2) and (3) and rearranging gives: DT0 σf − ln ε˙ = ln 3η0 T − T0
(3)
(4)
A plot of ln(˙ε) as a function 1/T − T0 is expected to yield straight, parallel lines of gradient DT0 and an intercept related to σ f : this explains the results given in Fig. 8. As may be expected, higher strains in the SCL region results in a deviation from linearity due to localized thinning of the samples, although it may also be a cumulative effect of the viscosity decreasing with the increasing strain rate as flow conditions change from Newtonian to non-Newtonian due to an increase in sample free volume. As the temperature is increased above TX , a marked deviation in the flow rate between samples becomes apparent (Fig. 8). For the sample loaded at 20 N, there is a marked decrease in the rate at which the strain rate is increasing compared to those samples under higher load. Such deviation in flow behaviour is shown by an inflexion in the temperature/strain-rate curve at ∼143.4 ◦ C. Although not as pronounced, the sample loaded at 40 N also shows similar behaviour with an inflexion at ∼144.5 ◦ C immediately before failure. In comparison to the case of the sample held at 50 N, this information further supports the theory of deformation-induced stabilization of the supercooled liquid, whereby the strengthening response of crystallization has been delayed or completely suppressed for samples under higher loads and higher rate of deformation. Here it can be appreciated that this deformation method shows potential of greater elongations before failure due to crystallization hardening. 4. Conclusions
Fig. 8. Diagram showing instantaneous strain rates as a function of temperature during constant load testing. Included in the diagram is the DSC profile of an unstrained sample for the same heating rate and temperature interval.
The tensile flow behaviour of a Ca65 Mg15 Zn20 bulk metallic glass (BMG) was investigated in the supercooled liquid (SCL) region for a range of strain rates (10−3 –10−4 s−1 ) and temperatures (105–120 ◦ C). It was found that:
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• the flow stress increases to a maximum followed by a decrease in stress to a low steady state value; • a maximum elongation of ∼600% is achieved when testing at 110 ◦ C for a strain rate of 5 × 10−4 s−1 ; • due to the onset of crystallization in the SCL region, there is an increase in stress which indicates that, at elevated temperatures, the overall strength of the semi-crystalline material is higher than that of the fully amorphous material and • the stress increase associated with crystallization varies with test temperature and strain rate, such that the undesirable crystallization reaction is more readily suppressed at lower temperatures and higher strain rates. The tensile flow behaviour of the BMG was also investigated under conditions of constant load (20–50 N) and constant heating rate (5 ◦ C/min). It was found that: • an increase in the deformation temperature in the SCL region results in a concomitant increase in strain rate of the material which is be explained by current deformation and viscosity theory (i.e. Eqs. (2) and (3)) and • hardening associated with the onset of crystallization is more readily suppressed at higher loads/deformation rates which allows higher tensile strains (>800%) to be achieved. Acknowledgement The authors would like to thank the Australian Research Council (ARC) for partial funding of this work via the ARC Centre of Excellence for Design in Light Metals (CEO561574).
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