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Crystallization and mechanical behavior of (Hf, Zr)–Ti–Cu–Ni–Al metallic glasses
Journal of Non-Crystalline Solids 317 (2003) 112–117 www.elsevier.com/locate/jnoncrysol
Crystallization and mechanical behavior of (Hf, Zr)–Ti–Cu–Ni–Al metallic glasses X. Gu a
a,*
, T. Jiao b, L.J. Kecskes c, R.H. Woodman c, C. Fan a, K.T. Ramesh b, T.C. Hufnagel a
Department of Materials Science and Engineering, Johns Hopkins University, Maryland Hall 102, 3400 North Charles Street, Baltimore, MD 21218, USA b Department of Mechanical Engineering, Johns Hopkins University, Maryland Hall 102, 3400 North Charles Street, Baltimore, MD 21218, USA c US Army Research Laboratory, Aberdeen Proving Ground, MD 21005, USA
Abstract We have prepared a series of glass-forming alloys of composition (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 and (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 with x ¼ 0–1. The substitution of Hf for Zr reduces the glass-forming ability of these alloys; while all compositions of (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 can be cast in bulk form, amorphous samples of the Hf-rich compositions (x > 0:8) of (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 can only be produced by rapid solidification. The alloys show similar crystallization behavior, although we observe an intermediate phase in the Hf-rich glasses (Al16 Hf6 Ni7 ) that is not observed in the Zr-rich glasses. The YoungÕs modulus and the flow strength under quasi-static uniaxial compression increase linearly with increasing Hf content. Under dynamic loading, the failure strength of all of the alloys tested decreases with increasing strain rate. The differences in mechanical behavior between the Zr- and Hf-rich glasses can be rationalized in terms of the higher melting point of Hf. The addition of Hf increases the average bond energy of the alloy, resulting in a higher modulus and increasing the activation energy for atomic motion by which plastic deformation occurs. Ó 2003 Elsevier Science B.V. All rights reserved. PACS: 61.43.Dq; 62.20.Dc; 81.05.Kf; 81.70.Pg
1. Introduction Multi-component Zr-based bulk metallic glasses have been widely investigated [1–3]. These alloys can be prepared at low cooling rates (below
*
Corresponding author. Tel.: +1-410 516 0462; fax: +1-410 516 5293. E-mail address: [email protected] (X. Gu).
103 K/s) and usually have wide supercooled regions (as large as 50 K). Plastic deformation of bulk metallic glasses is controlled predominantly by shear localization [4,5]. As a result, they usually exhibit very limited overall plasticity, which restricts most engineering applications of bulk metallic glasses as a structural material. However, shear localization leads to self-sharpening behavior during dynamic impact, which is important for making efficient anti-armor penetrators [6,7]. To
0022-3093/03/$ - see front matter Ó 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0022-3093(02)01990-7
X. Gu et al. / Journal of Non-Crystalline Solids 317 (2003) 112–117
113
improve the ballistic performance, it would be useful to increase the density of the material. Recently, tungsten fibers have been introduced into Zr-based glass-forming alloys to improve the mechanical properties and increase the density [7]. Further increases in the composite density may be achieved by a higher-density glass matrix. In the present work, we describe high-density metallic glasses produced by substituting Hf for Zr in two known glass-forming systems: Zr52:5 Cu17:9 Ni14:6 Al10 Ti5 [3] and Zrx57 Ti5 Cu20 Ni8 Al10 [8]. We produced two series of Hf-based bulk metallic glasses, with composition (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 (with x ¼ 0, 1/3, 1/2, 2/3, 1) and (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 (with x ¼ 0, 0.09, 0.18, 0.20, 0.35, 0.40, 0.50, 0.60, 0.74, 0.80, 1). With increasing Hf content, the density of these alloys increases by about 67% and 73%, respectively. The glass-forming ability of both alloys, however, is degraded by the substitution of Hf for Zr, and there are some slight differences in crystallization behavior. Under quasi-static compression, the YoungÕs modulus (E), the yield strength (ry ) and the fracture strength (rf ) all increase linearly with Hf content. This behavior is due to an increase in the average bond energy when Hf is substituted for Zr. The plastic strain to failure (ep ) is small and independent of composition. Under dynamic compression, the failure strength of all compositions tested decreases with increasing strain rate.
The density of the alloys was measured using ArchimedesÕs method, and averaged over several samples cut from different positions along each rod. Glass-transition and crystallization temperatures were measured by differential scanning calorimetry (DSC) (Perkin–Elmer Pyris 1) while a differential thermal analyzer (DTA) (Perkin–Elmer DTA 7) was used to measure the melting temperatures. X-ray diffraction (XRD) patterns were measured on a Philips diffractometer using Cu Ka radiation, and high resolution electron microscope (HRTEM) images were taken on a Philips CM300 FEG TEM operated at 300 kV. Quasi-static uniaxial compression tests were performed on a servo-hydraulic testing machine under displacement control with a strain rate of 104 s1 . The sample strain was calculated from the platen displacement, corrected for the compliance of the machine. The YoungÕs modulus was measured from the strain–stress curves. We measured the YoungÕs modulus of some standards and found the deviation of the calculated values is within 2%. Right circular cylinder samples 2.8 mm diameter and 5.6 mm long were machined from the rods for the quasi-static compression tests. The behavior under dynamic loading was studied using a compression Kolsky bar; the dynamic compression samples had a height-diameter ratio of about 0.6. The maximum strain rate was around 4000 s1 . All these tests were performed at room temperature.
2. Experimental
3. Results
Both (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 and (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 alloys were prepared by mixing pure elements (purity > 99:97%) according to their nominal weight ratios and arcmelting in Ti-gettered argon atmosphere (5 104 Pa). Each ingot was melted several times to improve homogeneity. For producing bulk samples, the arc-melted ingot was re-melted and cast into a rod-shaped cavity in a copper mold [9]. Ribbon samples were produced by single-roller melt spinning. We analyzed the oxygen contamination in several of the final ingots and found it to less than 100 ppm by weight.
Fig. 1 shows a typical HRTEM image and diffraction pattern obtained from the alloy (Hf0:5 Zr0:5 )52:5 Cu17:9 Ni14:6 Al10 Ti5 . Both XRD patterns [9] and HRTEM images indicate that the alloys are homogeneous and amorphous. From DSC and DTA experiments with the heating rate of 0.167 K/s we find that the glass transition temperature (Tg ), the crystallization temperatures (Tx ) and the melting temperature (solidus temperature, Tm ) increase linearly with increasing Hf content for both systems (Fig. 2(a)), while the reduced glass transition temperature (Trg ¼ Tg = Tm ) decreases with the substitution of Hf for Zr
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Fig. 1. A typical HREM image and electron diffraction pattern from the (Hf0:5 Zr0:5 )52:5 Cu17:9 Ni14:6 Al10 Ti5 glass.
(Fig. 2(b)). This indicates that the systematic substitution of Hf for Zr degrades the glass forming ability. The density increases linearly from 6:74 0:13 g/cm3 (x ¼ 0) to 11:08 0:17 g/cm3 (x ¼ 1) for (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 glasses and from 6:46 0:21 g/cm3 (x ¼ 0) to 11:17 0:19 g/cm3 (x ¼ 1) for (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 glasses.
For Zr52:5 Cu17:9 Ni14:6 Al10 Ti5 , constant heatingrate DSC scans show four exotherms during the crystallization process [9]. Isothermal DSC results (not shown) indicate that the first two exotherms represent nucleation and growth processes, while the third and the fourth exotherms have an exponential decay, corresponding to growth of existing phases. From this, and the XRD patterns shown in Fig. 3, we deduce that both Zr2 Ni and CuZr2 nucleate at around 700 K, but that the CuZr2 phase grows at a higher temperature (845 K) than the Zr2 Ni phase (780 K). With increasing annealing temperature and time, the fraction of the stable CuZr2 phase increases, while the Zr2 Ni phase decreases gradually (but does not totally disappear even after 8 h of isothermal annealing at 943 K). Applying a similar procedure to Hf52:5 Cu17:9 Ni14:6 Al10 Ti5 , we find that only the first exotherm corresponds to a nucleation and growth process. However, instead of Hf2 Ni as a counterpart of Zr2 Ni, a ternary intermediate phase Al16 Hf6 Ni7 grows first (848 K). At a higher temperature (918 K), the stable CuHf2 phase grows quickly with a slight decrease in the volume fraction of Al16 Hf6 Ni7 (Fig. 4). The isothermal annealing and XRD patterns of (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5
Fig. 2. (a) The glass transition temperature and melting temperature, and (b) and the reduced glass transition temperature of (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 (solid symbols) and (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 (open symbols) metallic glasses.
X. Gu et al. / Journal of Non-Crystalline Solids 317 (2003) 112–117
115
Fig. 3. X-ray diffraction patterns (Cu Ka radiation) of Zr52:5 Cu17:9 Ni14:6 Al10 Ti5 glasses that were annealed isothermally at the temperatures indicated for 90 min. The intensity axis is offset for clarity.
Fig. 4. X-ray diffraction patterns (Cu Ka radiation) of Hf52:5 Cu17:9 Ni14:6 Al10 Ti5 glasses that were annealed isothermally at the temperatures indicated for 90 min. The intensity axis is offset for clarity.
alloys give similar results as (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 alloys. Ni-containing phases grow first and more stable Cu-containing phases dominate finally. Fig. 5 shows the results of quasi-static compression tests on (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 . The YoungÕs modulus (E) increases linearly with Hf content from 86.6 GPa (x ¼ 0) to 102.4 GPa (x ¼ 1). The yield stress and the fracture stress
show similar behavior, reaching a maximum fracture stress of about 2150 MPa for Hf52:5 Cu17:9 Ni14:6 Al10 Ti5 . The overall plastic strain of these samples varies from 1.0% to 2.5%, but is not sensitive to the composition. From compression Kolsky bar tests (shown in Fig. 6), we find that the normalized failure stress (the dynamic failure stress divided by the quasistatic failure stress) decreases with increasing
Fig. 5. (a) The composition dependence of the YoungÕs modulus and (b) the yield stress and the fracture stress for (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 (x ¼ 0–1) glasses. The strain rate of quasi-static compression tests is 104 s1 .
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X. Gu et al. / Journal of Non-Crystalline Solids 317 (2003) 112–117
Fig. 6. Normalized failure stress versus strain rate for (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 (x ¼ 0–1), Zr41:2 Ti13:8 Cu12:5 Ni10 Be22:5 [10] and Zr57 Cu20 Ni8 Al10 Ti5 [11].
strain rate. Bruck et al. [10] measured the dynamic failure strength of beryllium-bearing Zr-based glasses; they observed a similar decrease, but attributed it to the effects of dispersion in their experiment. In the present case, our data are corrected for dispersion effects [11], and so we conclude that the observed trends are real.
4. Discussion Since Hf and Zr are chemically quite similar and have nearly the same atomic size, the density-composition relation can be easily understood. Hf replaces Zr in the structure of the glass, and the short-range order of the alloy is presumably affected little by the substitution. For this reason, partial substitution of Hf for Zr does not improve the glassforming ability of these alloys. This does not, however, exclude the possibility of forming different phases during annealing. Differences in the melting temperature of Hf and Zr, the bonding energies of the intermetallics, and the diffusion coefficients can change the characteristic temperatures of both cooperative and diffusion controlled processes, and produce different crystalline precipitates. The differences in mechanical properties due to Hf content can be understood in terms of the average bond energy in the alloy. In addition to its higher melting point, the heat of vaporization of
Hf is larger than that of Zr, indicative of a greater bond energy for Hf. Therefore, it is reasonable to suppose that the average bond energy in the amorphous solid increases with Hf content; this idea is supported by the observation that all of the characteristic temperatures (glass transition, crystallization, and melting) increase with Hf content. It is also confirmed by our observation of increased activation energy for glass transition and crystallization events evaluated by calorimetry (not shown). In the conventional theory of linear elasticity, YoungÕs modulus is related to the curvature of the interatomic potential energy well [12]. With increasing Hf content, the potential energy well for an average bond becomes deeper (reflecting the greater bond energy), but the interatomic spacing changes little (due to the similar atomic size of Zr and Hf). Thus, the curvature of the potential energy well (and, hence, the modulus) increases with Hf content. The increase in yield strength with increasing Hf content can be understood in a similar way. In any microscopic model of deformation, there will be a term representing an activation barrier for atomic motion. This activation barrier will be sensitive to the average bond energy. In the free volume model [13], for instance, there is an activation barrier associated with moving an atom into an adjacent jump site. In this case, the activation energy is due to the displacement of adjacent atoms as the moving atom makes its jump. It is reasonable to suppose that this activation energy scales with the average bond energy. Hence, the yield strength of the amorphous alloy also increases with increasing melting temperature. The observation that the dynamic failure strength of the amorphous alloys is lower than the quasi-static failure strength is apparently related to the dominant failure mechanism. Failure in these alloys is accompanied by significant local heating, which suggests that adiabatic processes occur during failure [11]. Since higher strain rates clearly favor adiabatic processes, this can explain why the failure stress is lower at the higher loading rates. It is interesting to note, however, that the failure strengths of all the alloys shown in Fig. 6 have essentially the same dependence on strain rate
X. Gu et al. / Journal of Non-Crystalline Solids 317 (2003) 112–117
(when normalized to the quasi-static failure strength), despite the fact that both the glass transition and melting temperatures are considerably higher in the Hf-rich alloys than in the Zr-rich alloys (Fig. 2).
117
University Center for Advanced Metallic and Ceramic Systems, under grant number DAAD19-012-0003; and by the National Science Foundation, under grant DMR-9875115.
References 5. Conclusions Two systems of (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 and (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 (x ¼ 0–1) bulk metallic glasses were successfully produced. The alloy density increases significantly when Zr is completely replaced by Hf, while the glass-forming ability decreases slightly. The exothermic DSC peaks and crystallization processes are investigated. The final stable crystalline phases are tetragonal CuZr2 and CuHf2 . The YoungÕs modulus, the yield strength and the fracture strength also increase with Hf content. The failure stress decreases with increasing strain rate, but shows no obvious composition dependence.
Acknowledgements This work was supported by the US Army Research Laboratory, through the Johns Hopkins
[1] A. Inoue, H. Yamaguchi, T. Zhang, T. Masumoto, Mater. Trans. JIM 31 (1990) 104. [2] A. Peker, W.L. Johnson, Appl. Phys. Lett. 63 (1993) 2342. [3] X.H. Lin, W.L. Johnson, W.K. Rhim, Mater. Trans. JIM 39 (1997) 477. [4] C.A. Pampillo, Scripta Metall. 6 (1972) 915; C.A. Pampillo, J. Mater. Sci. 10 (1975) 1194. [5] F. Spaepen, Acta Metall. 25 (1977) 407; J.C.M. Li, Metall. Trans. A 16A (1985) 2227. [6] L.S. Magness Jr., Mech. Mater. 17 (1994) 147. [7] C.-Y. Haein, R.D. Conner, F. Szuecs, W.L. Johnson, Scripta Mater. 45 (2001) 1039. [8] L.Q. Xing, J. Eckert, W. Loser, L. Schultz, Appl. Phys. Lett. 73 (1998) 2110. [9] X. Gu, L.Q. Xing, T.C. Hufnagel, J. Non-Cryst. Solids 311 (2002) 77. [10] H.A. Bruck, A.J. Rosakis, W.L. Johnson, J. Mater. Res. 11 (1996) 503. [11] T.C. Hufnagel, T. Jiao, Y. Li, L.-Q. Xing, K.T. Ramesh, J. Mater. Res. 17 (2002) 1441. [12] T.H. Courtney, Mechanical Behavior of Materials, McGraw-Hill, Boston, MA, 2000. [13] D. Turnbull, J.C. Fisher, J. Chem. Phys. 17 (1949) 71; D. Turnbull, J. Chem. Phys. 18 (1950) 198.
Crystallization and mechanical behavior of (Hf, Zr)–Ti–Cu–Ni–Al metallic glasses X. Gu a
a,*
, T. Jiao b, L.J. Kecskes c, R.H. Woodman c, C. Fan a, K.T. Ramesh b, T.C. Hufnagel a
Department of Materials Science and Engineering, Johns Hopkins University, Maryland Hall 102, 3400 North Charles Street, Baltimore, MD 21218, USA b Department of Mechanical Engineering, Johns Hopkins University, Maryland Hall 102, 3400 North Charles Street, Baltimore, MD 21218, USA c US Army Research Laboratory, Aberdeen Proving Ground, MD 21005, USA
Abstract We have prepared a series of glass-forming alloys of composition (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 and (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 with x ¼ 0–1. The substitution of Hf for Zr reduces the glass-forming ability of these alloys; while all compositions of (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 can be cast in bulk form, amorphous samples of the Hf-rich compositions (x > 0:8) of (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 can only be produced by rapid solidification. The alloys show similar crystallization behavior, although we observe an intermediate phase in the Hf-rich glasses (Al16 Hf6 Ni7 ) that is not observed in the Zr-rich glasses. The YoungÕs modulus and the flow strength under quasi-static uniaxial compression increase linearly with increasing Hf content. Under dynamic loading, the failure strength of all of the alloys tested decreases with increasing strain rate. The differences in mechanical behavior between the Zr- and Hf-rich glasses can be rationalized in terms of the higher melting point of Hf. The addition of Hf increases the average bond energy of the alloy, resulting in a higher modulus and increasing the activation energy for atomic motion by which plastic deformation occurs. Ó 2003 Elsevier Science B.V. All rights reserved. PACS: 61.43.Dq; 62.20.Dc; 81.05.Kf; 81.70.Pg
1. Introduction Multi-component Zr-based bulk metallic glasses have been widely investigated [1–3]. These alloys can be prepared at low cooling rates (below
*
Corresponding author. Tel.: +1-410 516 0462; fax: +1-410 516 5293. E-mail address: [email protected] (X. Gu).
103 K/s) and usually have wide supercooled regions (as large as 50 K). Plastic deformation of bulk metallic glasses is controlled predominantly by shear localization [4,5]. As a result, they usually exhibit very limited overall plasticity, which restricts most engineering applications of bulk metallic glasses as a structural material. However, shear localization leads to self-sharpening behavior during dynamic impact, which is important for making efficient anti-armor penetrators [6,7]. To
0022-3093/03/$ - see front matter Ó 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0022-3093(02)01990-7
X. Gu et al. / Journal of Non-Crystalline Solids 317 (2003) 112–117
113
improve the ballistic performance, it would be useful to increase the density of the material. Recently, tungsten fibers have been introduced into Zr-based glass-forming alloys to improve the mechanical properties and increase the density [7]. Further increases in the composite density may be achieved by a higher-density glass matrix. In the present work, we describe high-density metallic glasses produced by substituting Hf for Zr in two known glass-forming systems: Zr52:5 Cu17:9 Ni14:6 Al10 Ti5 [3] and Zrx57 Ti5 Cu20 Ni8 Al10 [8]. We produced two series of Hf-based bulk metallic glasses, with composition (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 (with x ¼ 0, 1/3, 1/2, 2/3, 1) and (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 (with x ¼ 0, 0.09, 0.18, 0.20, 0.35, 0.40, 0.50, 0.60, 0.74, 0.80, 1). With increasing Hf content, the density of these alloys increases by about 67% and 73%, respectively. The glass-forming ability of both alloys, however, is degraded by the substitution of Hf for Zr, and there are some slight differences in crystallization behavior. Under quasi-static compression, the YoungÕs modulus (E), the yield strength (ry ) and the fracture strength (rf ) all increase linearly with Hf content. This behavior is due to an increase in the average bond energy when Hf is substituted for Zr. The plastic strain to failure (ep ) is small and independent of composition. Under dynamic compression, the failure strength of all compositions tested decreases with increasing strain rate.
The density of the alloys was measured using ArchimedesÕs method, and averaged over several samples cut from different positions along each rod. Glass-transition and crystallization temperatures were measured by differential scanning calorimetry (DSC) (Perkin–Elmer Pyris 1) while a differential thermal analyzer (DTA) (Perkin–Elmer DTA 7) was used to measure the melting temperatures. X-ray diffraction (XRD) patterns were measured on a Philips diffractometer using Cu Ka radiation, and high resolution electron microscope (HRTEM) images were taken on a Philips CM300 FEG TEM operated at 300 kV. Quasi-static uniaxial compression tests were performed on a servo-hydraulic testing machine under displacement control with a strain rate of 104 s1 . The sample strain was calculated from the platen displacement, corrected for the compliance of the machine. The YoungÕs modulus was measured from the strain–stress curves. We measured the YoungÕs modulus of some standards and found the deviation of the calculated values is within 2%. Right circular cylinder samples 2.8 mm diameter and 5.6 mm long were machined from the rods for the quasi-static compression tests. The behavior under dynamic loading was studied using a compression Kolsky bar; the dynamic compression samples had a height-diameter ratio of about 0.6. The maximum strain rate was around 4000 s1 . All these tests were performed at room temperature.
2. Experimental
3. Results
Both (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 and (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 alloys were prepared by mixing pure elements (purity > 99:97%) according to their nominal weight ratios and arcmelting in Ti-gettered argon atmosphere (5 104 Pa). Each ingot was melted several times to improve homogeneity. For producing bulk samples, the arc-melted ingot was re-melted and cast into a rod-shaped cavity in a copper mold [9]. Ribbon samples were produced by single-roller melt spinning. We analyzed the oxygen contamination in several of the final ingots and found it to less than 100 ppm by weight.
Fig. 1 shows a typical HRTEM image and diffraction pattern obtained from the alloy (Hf0:5 Zr0:5 )52:5 Cu17:9 Ni14:6 Al10 Ti5 . Both XRD patterns [9] and HRTEM images indicate that the alloys are homogeneous and amorphous. From DSC and DTA experiments with the heating rate of 0.167 K/s we find that the glass transition temperature (Tg ), the crystallization temperatures (Tx ) and the melting temperature (solidus temperature, Tm ) increase linearly with increasing Hf content for both systems (Fig. 2(a)), while the reduced glass transition temperature (Trg ¼ Tg = Tm ) decreases with the substitution of Hf for Zr
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X. Gu et al. / Journal of Non-Crystalline Solids 317 (2003) 112–117
Fig. 1. A typical HREM image and electron diffraction pattern from the (Hf0:5 Zr0:5 )52:5 Cu17:9 Ni14:6 Al10 Ti5 glass.
(Fig. 2(b)). This indicates that the systematic substitution of Hf for Zr degrades the glass forming ability. The density increases linearly from 6:74 0:13 g/cm3 (x ¼ 0) to 11:08 0:17 g/cm3 (x ¼ 1) for (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 glasses and from 6:46 0:21 g/cm3 (x ¼ 0) to 11:17 0:19 g/cm3 (x ¼ 1) for (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 glasses.
For Zr52:5 Cu17:9 Ni14:6 Al10 Ti5 , constant heatingrate DSC scans show four exotherms during the crystallization process [9]. Isothermal DSC results (not shown) indicate that the first two exotherms represent nucleation and growth processes, while the third and the fourth exotherms have an exponential decay, corresponding to growth of existing phases. From this, and the XRD patterns shown in Fig. 3, we deduce that both Zr2 Ni and CuZr2 nucleate at around 700 K, but that the CuZr2 phase grows at a higher temperature (845 K) than the Zr2 Ni phase (780 K). With increasing annealing temperature and time, the fraction of the stable CuZr2 phase increases, while the Zr2 Ni phase decreases gradually (but does not totally disappear even after 8 h of isothermal annealing at 943 K). Applying a similar procedure to Hf52:5 Cu17:9 Ni14:6 Al10 Ti5 , we find that only the first exotherm corresponds to a nucleation and growth process. However, instead of Hf2 Ni as a counterpart of Zr2 Ni, a ternary intermediate phase Al16 Hf6 Ni7 grows first (848 K). At a higher temperature (918 K), the stable CuHf2 phase grows quickly with a slight decrease in the volume fraction of Al16 Hf6 Ni7 (Fig. 4). The isothermal annealing and XRD patterns of (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5
Fig. 2. (a) The glass transition temperature and melting temperature, and (b) and the reduced glass transition temperature of (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 (solid symbols) and (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 (open symbols) metallic glasses.
X. Gu et al. / Journal of Non-Crystalline Solids 317 (2003) 112–117
115
Fig. 3. X-ray diffraction patterns (Cu Ka radiation) of Zr52:5 Cu17:9 Ni14:6 Al10 Ti5 glasses that were annealed isothermally at the temperatures indicated for 90 min. The intensity axis is offset for clarity.
Fig. 4. X-ray diffraction patterns (Cu Ka radiation) of Hf52:5 Cu17:9 Ni14:6 Al10 Ti5 glasses that were annealed isothermally at the temperatures indicated for 90 min. The intensity axis is offset for clarity.
alloys give similar results as (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 alloys. Ni-containing phases grow first and more stable Cu-containing phases dominate finally. Fig. 5 shows the results of quasi-static compression tests on (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 . The YoungÕs modulus (E) increases linearly with Hf content from 86.6 GPa (x ¼ 0) to 102.4 GPa (x ¼ 1). The yield stress and the fracture stress
show similar behavior, reaching a maximum fracture stress of about 2150 MPa for Hf52:5 Cu17:9 Ni14:6 Al10 Ti5 . The overall plastic strain of these samples varies from 1.0% to 2.5%, but is not sensitive to the composition. From compression Kolsky bar tests (shown in Fig. 6), we find that the normalized failure stress (the dynamic failure stress divided by the quasistatic failure stress) decreases with increasing
Fig. 5. (a) The composition dependence of the YoungÕs modulus and (b) the yield stress and the fracture stress for (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 (x ¼ 0–1) glasses. The strain rate of quasi-static compression tests is 104 s1 .
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Fig. 6. Normalized failure stress versus strain rate for (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 (x ¼ 0–1), Zr41:2 Ti13:8 Cu12:5 Ni10 Be22:5 [10] and Zr57 Cu20 Ni8 Al10 Ti5 [11].
strain rate. Bruck et al. [10] measured the dynamic failure strength of beryllium-bearing Zr-based glasses; they observed a similar decrease, but attributed it to the effects of dispersion in their experiment. In the present case, our data are corrected for dispersion effects [11], and so we conclude that the observed trends are real.
4. Discussion Since Hf and Zr are chemically quite similar and have nearly the same atomic size, the density-composition relation can be easily understood. Hf replaces Zr in the structure of the glass, and the short-range order of the alloy is presumably affected little by the substitution. For this reason, partial substitution of Hf for Zr does not improve the glassforming ability of these alloys. This does not, however, exclude the possibility of forming different phases during annealing. Differences in the melting temperature of Hf and Zr, the bonding energies of the intermetallics, and the diffusion coefficients can change the characteristic temperatures of both cooperative and diffusion controlled processes, and produce different crystalline precipitates. The differences in mechanical properties due to Hf content can be understood in terms of the average bond energy in the alloy. In addition to its higher melting point, the heat of vaporization of
Hf is larger than that of Zr, indicative of a greater bond energy for Hf. Therefore, it is reasonable to suppose that the average bond energy in the amorphous solid increases with Hf content; this idea is supported by the observation that all of the characteristic temperatures (glass transition, crystallization, and melting) increase with Hf content. It is also confirmed by our observation of increased activation energy for glass transition and crystallization events evaluated by calorimetry (not shown). In the conventional theory of linear elasticity, YoungÕs modulus is related to the curvature of the interatomic potential energy well [12]. With increasing Hf content, the potential energy well for an average bond becomes deeper (reflecting the greater bond energy), but the interatomic spacing changes little (due to the similar atomic size of Zr and Hf). Thus, the curvature of the potential energy well (and, hence, the modulus) increases with Hf content. The increase in yield strength with increasing Hf content can be understood in a similar way. In any microscopic model of deformation, there will be a term representing an activation barrier for atomic motion. This activation barrier will be sensitive to the average bond energy. In the free volume model [13], for instance, there is an activation barrier associated with moving an atom into an adjacent jump site. In this case, the activation energy is due to the displacement of adjacent atoms as the moving atom makes its jump. It is reasonable to suppose that this activation energy scales with the average bond energy. Hence, the yield strength of the amorphous alloy also increases with increasing melting temperature. The observation that the dynamic failure strength of the amorphous alloys is lower than the quasi-static failure strength is apparently related to the dominant failure mechanism. Failure in these alloys is accompanied by significant local heating, which suggests that adiabatic processes occur during failure [11]. Since higher strain rates clearly favor adiabatic processes, this can explain why the failure stress is lower at the higher loading rates. It is interesting to note, however, that the failure strengths of all the alloys shown in Fig. 6 have essentially the same dependence on strain rate
X. Gu et al. / Journal of Non-Crystalline Solids 317 (2003) 112–117
(when normalized to the quasi-static failure strength), despite the fact that both the glass transition and melting temperatures are considerably higher in the Hf-rich alloys than in the Zr-rich alloys (Fig. 2).
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University Center for Advanced Metallic and Ceramic Systems, under grant number DAAD19-012-0003; and by the National Science Foundation, under grant DMR-9875115.
References 5. Conclusions Two systems of (Hfx Zr1x )52:5 Cu17:9 Ni14:6 Al10 Ti5 and (Hfx Zr1x )57 Cu20 Ni8 Al10 Ti5 (x ¼ 0–1) bulk metallic glasses were successfully produced. The alloy density increases significantly when Zr is completely replaced by Hf, while the glass-forming ability decreases slightly. The exothermic DSC peaks and crystallization processes are investigated. The final stable crystalline phases are tetragonal CuZr2 and CuHf2 . The YoungÕs modulus, the yield strength and the fracture strength also increase with Hf content. The failure stress decreases with increasing strain rate, but shows no obvious composition dependence.
Acknowledgements This work was supported by the US Army Research Laboratory, through the Johns Hopkins
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