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Effects of partial replacement of Cu and Y by Bin Mg–Cu–Y amorphous alloys
Intermetallics 14 (2006) 866–870 www.elsevier.com/locate/intermet
Effects of partial replacement of Cu and Y by B in Mg–Cu–Y amorphous alloys Y.T. Cheng a, T.H. Hung a, J.C. Huang a,*, J.S.C. Jang b, Chi Y.A. Tsao c, P.Y. Lee d a
Institute of Materials Science and Engineering; Center for Nanoscience and Nanotechnology, National Sun Yat-Sen University, Kaohsiung, Taiwan, ROC b Department of Materials Science and Engineering, I-Shou University, Kaohsiung, Taiwan 804, ROC c Department of Materials Science and Engineering, National Cheng Kung University, Tainan, Taiwan, ROC d Institute of Materials Engineering, National Taiwan Ocean University, Keelung 202, Taiwan, ROC Available online 20 February 2006
Abstract The effects of replacing Cu or Y by a small-sized B in the Mg65Cu25Y10 based alloy are examined. It is found that the replacement of Y by B consistently leads to apparent degradation in GFA; the large sized Y seems to be irreplaceable. In contrast, the replacement of the small sized Cu by the even smaller B to a small amount (1–5%) appears to be beneficial in terms of wider DT, higher GFA, thermal stability, and hardness, as well as lower density. In this study, the optimum quaternary alloy composition is identified to be Mg65Cu22Y10B3. q 2006 Elsevier Ltd. All rights reserved. Keywords: B. Glasses, metallic; B. Thermal stability; C. Rapid solidification processing; F. Calorimetry
1. Introduction Among a large number of amorphous alloys, Mg-based alloys have attracted much attention especially due to its high strength to weight ratio and a low glass transition temperature. The new Mg-based amorphous alloys with high tensile strength and good ductility were first found in 1988 [1]. In 1991, Inoue et al. [2] were succeessful in finding new Mg-based amorphous alloys, such as Mg–Ln–TM (LnZlanthanide metal; TMZ transition metal) systems with high tensile strength. They also fabricated bulk metallic glass (BMG) with a diameter of 4 mm [3] by injection casting the Mg65Cu25Y10 alloy into a Cu mold. The next year, Inoue’s group used high-pressure die casting method and increased the diameter even more to 7 mm [4]. Their contribution opened the door to design new families of light amorphous alloy systems. Inoue et al. [5] have reported the relationship between the thermal stability of an amorphous phase and the topological short range ordering. It is thought that the increase in packing density can cause an increase in thermal stability. The follow-up investigations on the Mg-based BMG have been directed to the addition of the quaternary element to improve the glass forming ability (GFA). Recently, Park et * Corresponding author. Tel.: C886 7525 2000; fax: C886 7 525 4099. E-mail address: [email protected] (J.C. Huang).
0966-9795/$ - see front matter q 2006 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2006.01.002
al. [6] reported the effect of Ag as the substituting element for Cu. Due to the combined effects of a large difference in atomic size between Ag and the constituting elements and the larger mixing enthalpy of Ag–Mg and Ag–Y than those of Cu, the addition of Ag would contribute to increase the atomic packing density of the liquid structure and GFA. Men et al. [7] also investigated the effect of substitution of Y in the Mg65Cu25Y10 alloy with Gd on the GFA and crystallization behavior. Due to the difference in electronic configuration between Y (4d15s2) and Gd (4f75d16s2), it results in a certain change of the short-range order in the supercooled liquid of Mg–Cu–Y–Gd alloys and shows a higher GFA. In addition, it has been reported that the Mg65Cu20Zn5Y10 [8] and Mg65Cu15Ag5Pd5Y10 [9] alloys also exhibit a high GFA and enable to fabricate metallic rods with diameters of 6 and 7 mm by a Cu-mold injection casting. From the above review, it seems that a proper addition of some quaternary elements might play an important role in improving the GFA of the original Mg65Cu25Y10 amorphous alloys. In this paper, we studied the addition of the interstitial atom, B, into Mg65Cu25Y10. It could cause a large difference in atomic size in the Mg–Cu–Y–B system and would also induce strongly negative DHm for Mg-B and Y–B (both around K50 kJ/mol). Also, the addition of B might fill the liquid free volume and increase the density of randomly packed liquids; both would further improve the GFA and thermal stability against crystallization. Furthermore, boron has a lower density
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Fig. 1. XRD patterns for the (a) Mg65Cu25Y10KxBx and (b) Mg65Cu25KxY10Bx systems.
of 2.35 g/cm3 than that of Cu and Y thus, the resulting quaternary alloy might possess a lower density, which would be helpful to the lightweight and impose no harm on GFA. 2. Experimental methods The master ingot of Cu–Y binary alloys was first prepared by arc melting high purity Cu (99.999%) and Y (99.9%) under a Tigettered argon atmosphere. Then the master ingot was melted with pure Mg (99.99%) and B (99.5%) in an induction furnace under a purified argon atmosphere. After complete melting, the liquid alloy was poured onto the surface of a single Cu wheel. The wheel rotated at a high speed of 25 m/s (20 Hz) in order to reach a high cooling rate, so as the glassy ribbon samples were rapidly solidified. The resulting specimen is thin ribbon type, w10 mm in width and w100 mm in thickness. In Mg–Cu–Y–B amorphous alloys fabrication, since the boiling point of Mg is much lower than the melting point of Cu, Y and B. Therefore, during the melting process, the crucible is designed to have a cap and argon atmosphere is added to 2 atm in order to minimize the evaporation of the pure Mg component. The vitrification degree of the ribbon structure is first identified by the SIEMENS D5000 X-ray diffractometer (XRD) using the Cu Ka radiation. Thermal analysis was carried out to determine the glass transition temperature Tg, crystallization
Fig. 2. Variation of the microhardness HV values as a function of B content for the Mg65Cu25KXY10BX alloys.
temperature Tx, and melting temperature Tm by the differential scanning calorimetry (DSC) under flowing argon atmosphere from 303 to 800 K. Heating rates of 10–40 K/min were employed to reveal the crystallization and melting behavior. The Vickers hardness tests were conducted using a 500 gf load for 10 s and the plastic deformation and shear bands around the intents in the amorphous alloys were examined using the scanning electron microscopy (SEM). 3. Results and discussions 3.1. Microstructures The Mg–Cu–Y metallic glass performs ductility behavior with 1508 bending angle. With increasing B content, the samples appear gradually more brittle and eventually become easily splintered into pieces as B reaches 10%. Fig. 1(a) shows the X-ray diffraction patterns of the melt spun Mg65Cu25Y10KxBx (xZ0, 1, 3, 5) alloys with B replacing Y. It is apparent that the Mg65Cu25Y10KxBx alloys exhibit an almost fully amorphous phase except a couple of small peaks superimposed on the amorphous diffused peak is observed at 2qw358 and 378 in the alloy with 5 at% boron content. The tiny peaks are identified to be the Mg2Cu crystalline phase with an orthorhombic structure. For a higher B addition of 10%, or completely substituting the yttrium, only the sharp Mg2Cu
Fig. 3. SEM micrograph showing the shear bands in the Mg65Cu25KXY10Bx alloy induced by the indentation.
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Fig. 4. DSC curves of the various amorphous alloys heated at a heating rate of 20 K/min: (a) Mg65Cu25KxY10Bx and (b) Mg65Cu25Y10-xBx.
crystalline peaks, together with some minor MgB4 crystalline peaks, are visible in the XRD curve. In other words, the Mg65Cu25B10 is no longer amorphous. On the other hand, from Fig. 1(b) for the melt spun Mg65Cu25KxY10Bx alloys with Cu being partially replaced by B, only the broad and diffused amorphous peak occurs and no obvious crystalline peak is seen in the melt spun Mg65Cu25KxY10Bx (xZ0, 1, 3, 5, 10) alloys. The XRD results imply that the partial substitution of Cu and Y by B has different effects on the GFA; replacing Y by B seems to impose negative effects, while replacing Cu by B results in a high GFA. 3.2. Hardness measurement Microhardness HV results of the Mg65Cu25KxY10Bx (with xZ0–10%) systems are shown in Fig. 2. With the addition of B, the HV increases appreciably from 220 to more than 300, a nearly 50% increment. Coupled with the lower density of the Mg65Cu25KxY10Bx alloys by around 5–15% as compared with the parent Mg65Cu25Y10, the new alloy possesses much higher specific mechanical properties. When the indenter impresses into the amorphous specimen, the plastic deformation morphology can be seen in the form of the pile-up of the semi-circular shear bands, as shown in Fig. 3.
amorphous phase and no Tg can be detected from its DSC curve (not shown in Fig. 4(b)). Meanwhile, the area of the exothermic peak of the Mg65Cu25KxY10Bx systems is much larger than that of Mg65Cu25Y10KxBx, implying that the higher crystallization enthalpy DHx and the higher packing density of the supercooled liquid in the former alloys. The variations of Tg and Tx as a function of B content for the Mg65Cu25KxY10Bx systems are presented in Fig. 5. The melting behavior of the Mg-based alloys is also determined using the DSC during continuous heating at a heating rate of 20 K/min. The reduced glass transition temperature, Trg, or the ratio of the glass transition temperature and the offset melting temperature (liquidus) Tl is often used as another parameter to estimate GFA. From the former study [10], Mg65Cu25Y10 would exhibit a single endothermic peak with a narrow melting range of about 40 K. The onset and offset temperatures of the melting at a heating rate of 20 K/min, as designated by Tmsolid and Tmliquid , are 728 and 770 K, respectively. It indicates that the Mg65Cu25Y10 alloy is close to the ternary eutectic composition. Compared with the parent Mg65Cu25Y10 alloy, from Fig. 6, Tmsolid of the xZ3 and 5 alloys for the Mg65Cu25KxY10Bx system decreases to about 715 and 712 K, and Tmliquid also decreases to about 738 and 736 K. Although the xZ3 alloy has two endothermic peaks, the lower melting point (736 K) and
3.3. Thermal analyses In the Mg65Cu25KxY10Bx systems, all the samples exhibit a distinct glass transition, followed by a wide supercooled region and then by exothermic reactions due to crystallization. Fig. 4(a) shows that, with increasing x from 0 to 3, the glass transition temperature is nearly constant at 410 K but the crystallization temperature increases slightly from 468 to 476 K. Therefore, the DTx value also slightly increases from 58 K at xZ0 to 66 K at xZ3, or there to a 14% increment in the DTx value. With further increasing x to 5 and 10 at%, the DTx value becomes smaller to 50 K for 5 at% B and 45 K for 10 at%. In contrast, DTx becomes progressively reduced in the Mg65Cu25Y10KxBx systems, as shown in Fig. 4(b). The Mg65Cu25B10 alloy contains no
Fig. 5. Variations of the glass transition and crystallization temperature as a function of B content for the Mg65Cu25KxY10Bx alloys.
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Furthermore, from the DSC trace curves, it is evident that the curve shape between Mg65Cu25KxY10Bx and Mg65Cu25Y10KxBx is greatly different. For instance, only one sharp exothermic peak with a high enthalpy of transformation is observed from the Mg65Cu22Y10B3 alloy as shown in Fig. 4(a). This might be a result that this alloy is presumed to have a higher degree of dense, random packed structure. In contrast, the Mg65Cu25Y7B3 alloy has two small exothermic peaks. In non-isothermal analyses, the dependence of crystallization temperature on heating rate could be used to estimate the associated crystallization activation energy by means of the Kissinger peak shift method [12], Fig. 6. DSC traces showing the melting endotherms obtained from the Mg65Cu25KxY10 Bx (xZ3, 5) alloys at a heating rate of 20 K/min.
Table 1 Thermal properties of the Mg65Cu25-xY10Bx (xZ0, 3, 5) alloys Compositions
Mg65Cu25Y10 Mg65Cu22Y10B3 Mg65Cu20Y10B5
Tg (K)
Tx (K)
DTx (K)
Tmsolid (K)
Tmliquid (K)
Tmliquid
410 410 420
468 476 470
58 66 50
728 715 712
770 738 736
0.532 0.556 0.571
Tg =
g
0.40 0.42 0.41
even smaller melting interval (23 K) indicate that the replacement of Cu by B would result in the apparent narrowing of solid/liquid two phase region and would enhance the GFA of the Mg–Cu–Y amorphous alloys. Also, the single endothermic peak and the narrow melting interval of the Mg65Cu25KxY10Bx with xZ5 demonstrate that this alloy may be very close to the quaternary eutectic composition. The Trg of the optimum Mg65Cu22Y10B3 alloy is 0.556, higher than that for the parent Mg65Cu25Y10 or it corresponds to a 5% increment. The summary of Tg, Tx, DTx, Tm and g (defined as Tx/TgC Tl) [11] is shown in Table 1. From the glass forming g index, the Mg65Cu22Y10B3 alloy again has the highest value of 0.42 than all others, or a 5% increment. Also, as a result of the narrow interval between Tm and Tx, it can contribute to the bulk amorphous formation without a high cooling rate.
lnðb=T 2 Þ ZKEa =RT C constant;
(1)
where b is the heating rate, T is the specific onset temperature, R is gas constant, and Ea is the associated activation energy. Fig. 7 shows the Kissinger’s plots of ln(b/Tx) against 1/Tx, taken from the dependence of crystallization temperature Tx on heating rate b (10, 20, 30 K/min) in the DSC curves. The associated activation energy of the parent Mg65Cu25Y10 amorphous alloy determined from the slope of Kissinger’s plot is 138 kJ/mol, as shown in Fig. 7(a). As a comparison, the activation energy for Mg65Cu22Y10B3 is calculated to be 156 kJ/mol, as compared in Fig. 7(b). The value is about 13% higher than the based alloy. It implies that the substitution of Cu by B would improve the thermal stability and impose a higher energy barrier against crystallization for the Mg-based alloys. The improvement of GFA in the Mg65Cu25KxY10Bx alloys might be a result of multiple factors, including the increment of packing density in supercooled liquid, the larger overall atomic size difference and the higher overall electronegativity difference [13]. The strong affinity for B/Mg and B/Y will suppress the formation of Mg2Cu, MgY and CuY phases, and indirectly enhance the thermal stability of the Mg65Cu25KxY10Bx alloys. Furthermore, both Y and B have a stronger affinity for oxygen than Mg and Cu. A small amount of B addition might be able to suppress the formation of yttrium oxide nuclei, but excessive B would again induce the formation of the B2O3 and YB4. It follows that a critical amount of B addition (w3%) appears to be optimum. A similar effect might also occur in the
Fig. 7. Kissinger plots for the melt spun alloys: (a) Mg65Cu25Y10 (b) Mg65Cu22Y10B3.
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Mg65Cu25Tb10 amorphous alloy [14], where Tb also shows high formation enthalpy to form TbO2. 4. Conclusions The effects of adding interstitial B atoms to the well-known Mg65Cu25Y10 amorphous alloy are examined. The modified Mg65Cu22Y10B3 alloy, with partial replacement of Cu by 3 at% B, exhibits a number of improved properties as compared to Mg65Cu25Y10, for example, a wider supercooled region of 66 K (14% increment versus the original 55 K for Mg65Cu25Y10), a narrower solid/liquid two-phase region of 23 K (43% reduction versus the original 40 K), a higher glass forming g index of 0.42 (5% increment versus the original 0.40), a higher reduced glass transition temperature Trg of 0.556 (5% increment versus the original 0.532), a larger crystallization activation energy of 156 kJ/mol (13% increment versus the original 138 kJ/mol) and an improved hardness of 310 HV (41% increment versus the original 220 HV). This demonastrated that by a proper addition of boron to the parent Mg65Cu25Y10 amorphous alloy, Mg-based amorphous alloys with better thermal stability and mechanical property can be developed.
Acknowledgements The authors gratefully acknowledge the sponsorship by the National Science Council of Taiwan, ROC, under the project no. NSC 93-2216-E-110-021. References [1] Inoue A, Ohtera K, Kita K, Masumoto T. Jpn J Appl Phys 1988;27:2248. [2] Inoue A, Masumoto T. Mater Sci Eng A 1991;133:6. [3] Inoue A, Kato A, Zhang T, Kim SG, Masumoto T. Mater Trans JIM 1991; 32:609. [4] Inoue A, Nakamura T, Nishiyama N, Masumoto T. Mater Trans JIM 1992;33:937. [5] Inoue A. Acta Mater 2000;48:279. [6] Park ES, Kang HG, Kim WT, Kim DH. J Non Cryst Solids 2001;279:154. [7] Men H, Kim WT, Kim DH. J Non Cryst Solids 2004;337:29. [8] Men H, Hu ZQ, Xu J. Scripta Mater 2002;46:699. [9] Amiya K, Inoue A. Mater Trans JIM 2000;41:1460. [10] Lu ZP, Li Y, Ng SC. J Non Cryst Solids 2000;270:103. [11] Lu ZP, Liu CT. Acta Mater 2002;50:3501. [12] Kissinger HE. Analyst Chem 1957;29:1702. [13] Fang S, Xiao X, Xia L, Li W, Dong Y. J Non Cryst Solids 2003;321:120. [14] Xi XK, Zhao DQ, Pan MX, Wang WH. J Non Cryst Solids 2004;344:189.
Effects of partial replacement of Cu and Y by B in Mg–Cu–Y amorphous alloys Y.T. Cheng a, T.H. Hung a, J.C. Huang a,*, J.S.C. Jang b, Chi Y.A. Tsao c, P.Y. Lee d a
Institute of Materials Science and Engineering; Center for Nanoscience and Nanotechnology, National Sun Yat-Sen University, Kaohsiung, Taiwan, ROC b Department of Materials Science and Engineering, I-Shou University, Kaohsiung, Taiwan 804, ROC c Department of Materials Science and Engineering, National Cheng Kung University, Tainan, Taiwan, ROC d Institute of Materials Engineering, National Taiwan Ocean University, Keelung 202, Taiwan, ROC Available online 20 February 2006
Abstract The effects of replacing Cu or Y by a small-sized B in the Mg65Cu25Y10 based alloy are examined. It is found that the replacement of Y by B consistently leads to apparent degradation in GFA; the large sized Y seems to be irreplaceable. In contrast, the replacement of the small sized Cu by the even smaller B to a small amount (1–5%) appears to be beneficial in terms of wider DT, higher GFA, thermal stability, and hardness, as well as lower density. In this study, the optimum quaternary alloy composition is identified to be Mg65Cu22Y10B3. q 2006 Elsevier Ltd. All rights reserved. Keywords: B. Glasses, metallic; B. Thermal stability; C. Rapid solidification processing; F. Calorimetry
1. Introduction Among a large number of amorphous alloys, Mg-based alloys have attracted much attention especially due to its high strength to weight ratio and a low glass transition temperature. The new Mg-based amorphous alloys with high tensile strength and good ductility were first found in 1988 [1]. In 1991, Inoue et al. [2] were succeessful in finding new Mg-based amorphous alloys, such as Mg–Ln–TM (LnZlanthanide metal; TMZ transition metal) systems with high tensile strength. They also fabricated bulk metallic glass (BMG) with a diameter of 4 mm [3] by injection casting the Mg65Cu25Y10 alloy into a Cu mold. The next year, Inoue’s group used high-pressure die casting method and increased the diameter even more to 7 mm [4]. Their contribution opened the door to design new families of light amorphous alloy systems. Inoue et al. [5] have reported the relationship between the thermal stability of an amorphous phase and the topological short range ordering. It is thought that the increase in packing density can cause an increase in thermal stability. The follow-up investigations on the Mg-based BMG have been directed to the addition of the quaternary element to improve the glass forming ability (GFA). Recently, Park et * Corresponding author. Tel.: C886 7525 2000; fax: C886 7 525 4099. E-mail address: [email protected] (J.C. Huang).
0966-9795/$ - see front matter q 2006 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2006.01.002
al. [6] reported the effect of Ag as the substituting element for Cu. Due to the combined effects of a large difference in atomic size between Ag and the constituting elements and the larger mixing enthalpy of Ag–Mg and Ag–Y than those of Cu, the addition of Ag would contribute to increase the atomic packing density of the liquid structure and GFA. Men et al. [7] also investigated the effect of substitution of Y in the Mg65Cu25Y10 alloy with Gd on the GFA and crystallization behavior. Due to the difference in electronic configuration between Y (4d15s2) and Gd (4f75d16s2), it results in a certain change of the short-range order in the supercooled liquid of Mg–Cu–Y–Gd alloys and shows a higher GFA. In addition, it has been reported that the Mg65Cu20Zn5Y10 [8] and Mg65Cu15Ag5Pd5Y10 [9] alloys also exhibit a high GFA and enable to fabricate metallic rods with diameters of 6 and 7 mm by a Cu-mold injection casting. From the above review, it seems that a proper addition of some quaternary elements might play an important role in improving the GFA of the original Mg65Cu25Y10 amorphous alloys. In this paper, we studied the addition of the interstitial atom, B, into Mg65Cu25Y10. It could cause a large difference in atomic size in the Mg–Cu–Y–B system and would also induce strongly negative DHm for Mg-B and Y–B (both around K50 kJ/mol). Also, the addition of B might fill the liquid free volume and increase the density of randomly packed liquids; both would further improve the GFA and thermal stability against crystallization. Furthermore, boron has a lower density
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Fig. 1. XRD patterns for the (a) Mg65Cu25Y10KxBx and (b) Mg65Cu25KxY10Bx systems.
of 2.35 g/cm3 than that of Cu and Y thus, the resulting quaternary alloy might possess a lower density, which would be helpful to the lightweight and impose no harm on GFA. 2. Experimental methods The master ingot of Cu–Y binary alloys was first prepared by arc melting high purity Cu (99.999%) and Y (99.9%) under a Tigettered argon atmosphere. Then the master ingot was melted with pure Mg (99.99%) and B (99.5%) in an induction furnace under a purified argon atmosphere. After complete melting, the liquid alloy was poured onto the surface of a single Cu wheel. The wheel rotated at a high speed of 25 m/s (20 Hz) in order to reach a high cooling rate, so as the glassy ribbon samples were rapidly solidified. The resulting specimen is thin ribbon type, w10 mm in width and w100 mm in thickness. In Mg–Cu–Y–B amorphous alloys fabrication, since the boiling point of Mg is much lower than the melting point of Cu, Y and B. Therefore, during the melting process, the crucible is designed to have a cap and argon atmosphere is added to 2 atm in order to minimize the evaporation of the pure Mg component. The vitrification degree of the ribbon structure is first identified by the SIEMENS D5000 X-ray diffractometer (XRD) using the Cu Ka radiation. Thermal analysis was carried out to determine the glass transition temperature Tg, crystallization
Fig. 2. Variation of the microhardness HV values as a function of B content for the Mg65Cu25KXY10BX alloys.
temperature Tx, and melting temperature Tm by the differential scanning calorimetry (DSC) under flowing argon atmosphere from 303 to 800 K. Heating rates of 10–40 K/min were employed to reveal the crystallization and melting behavior. The Vickers hardness tests were conducted using a 500 gf load for 10 s and the plastic deformation and shear bands around the intents in the amorphous alloys were examined using the scanning electron microscopy (SEM). 3. Results and discussions 3.1. Microstructures The Mg–Cu–Y metallic glass performs ductility behavior with 1508 bending angle. With increasing B content, the samples appear gradually more brittle and eventually become easily splintered into pieces as B reaches 10%. Fig. 1(a) shows the X-ray diffraction patterns of the melt spun Mg65Cu25Y10KxBx (xZ0, 1, 3, 5) alloys with B replacing Y. It is apparent that the Mg65Cu25Y10KxBx alloys exhibit an almost fully amorphous phase except a couple of small peaks superimposed on the amorphous diffused peak is observed at 2qw358 and 378 in the alloy with 5 at% boron content. The tiny peaks are identified to be the Mg2Cu crystalline phase with an orthorhombic structure. For a higher B addition of 10%, or completely substituting the yttrium, only the sharp Mg2Cu
Fig. 3. SEM micrograph showing the shear bands in the Mg65Cu25KXY10Bx alloy induced by the indentation.
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Fig. 4. DSC curves of the various amorphous alloys heated at a heating rate of 20 K/min: (a) Mg65Cu25KxY10Bx and (b) Mg65Cu25Y10-xBx.
crystalline peaks, together with some minor MgB4 crystalline peaks, are visible in the XRD curve. In other words, the Mg65Cu25B10 is no longer amorphous. On the other hand, from Fig. 1(b) for the melt spun Mg65Cu25KxY10Bx alloys with Cu being partially replaced by B, only the broad and diffused amorphous peak occurs and no obvious crystalline peak is seen in the melt spun Mg65Cu25KxY10Bx (xZ0, 1, 3, 5, 10) alloys. The XRD results imply that the partial substitution of Cu and Y by B has different effects on the GFA; replacing Y by B seems to impose negative effects, while replacing Cu by B results in a high GFA. 3.2. Hardness measurement Microhardness HV results of the Mg65Cu25KxY10Bx (with xZ0–10%) systems are shown in Fig. 2. With the addition of B, the HV increases appreciably from 220 to more than 300, a nearly 50% increment. Coupled with the lower density of the Mg65Cu25KxY10Bx alloys by around 5–15% as compared with the parent Mg65Cu25Y10, the new alloy possesses much higher specific mechanical properties. When the indenter impresses into the amorphous specimen, the plastic deformation morphology can be seen in the form of the pile-up of the semi-circular shear bands, as shown in Fig. 3.
amorphous phase and no Tg can be detected from its DSC curve (not shown in Fig. 4(b)). Meanwhile, the area of the exothermic peak of the Mg65Cu25KxY10Bx systems is much larger than that of Mg65Cu25Y10KxBx, implying that the higher crystallization enthalpy DHx and the higher packing density of the supercooled liquid in the former alloys. The variations of Tg and Tx as a function of B content for the Mg65Cu25KxY10Bx systems are presented in Fig. 5. The melting behavior of the Mg-based alloys is also determined using the DSC during continuous heating at a heating rate of 20 K/min. The reduced glass transition temperature, Trg, or the ratio of the glass transition temperature and the offset melting temperature (liquidus) Tl is often used as another parameter to estimate GFA. From the former study [10], Mg65Cu25Y10 would exhibit a single endothermic peak with a narrow melting range of about 40 K. The onset and offset temperatures of the melting at a heating rate of 20 K/min, as designated by Tmsolid and Tmliquid , are 728 and 770 K, respectively. It indicates that the Mg65Cu25Y10 alloy is close to the ternary eutectic composition. Compared with the parent Mg65Cu25Y10 alloy, from Fig. 6, Tmsolid of the xZ3 and 5 alloys for the Mg65Cu25KxY10Bx system decreases to about 715 and 712 K, and Tmliquid also decreases to about 738 and 736 K. Although the xZ3 alloy has two endothermic peaks, the lower melting point (736 K) and
3.3. Thermal analyses In the Mg65Cu25KxY10Bx systems, all the samples exhibit a distinct glass transition, followed by a wide supercooled region and then by exothermic reactions due to crystallization. Fig. 4(a) shows that, with increasing x from 0 to 3, the glass transition temperature is nearly constant at 410 K but the crystallization temperature increases slightly from 468 to 476 K. Therefore, the DTx value also slightly increases from 58 K at xZ0 to 66 K at xZ3, or there to a 14% increment in the DTx value. With further increasing x to 5 and 10 at%, the DTx value becomes smaller to 50 K for 5 at% B and 45 K for 10 at%. In contrast, DTx becomes progressively reduced in the Mg65Cu25Y10KxBx systems, as shown in Fig. 4(b). The Mg65Cu25B10 alloy contains no
Fig. 5. Variations of the glass transition and crystallization temperature as a function of B content for the Mg65Cu25KxY10Bx alloys.
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Furthermore, from the DSC trace curves, it is evident that the curve shape between Mg65Cu25KxY10Bx and Mg65Cu25Y10KxBx is greatly different. For instance, only one sharp exothermic peak with a high enthalpy of transformation is observed from the Mg65Cu22Y10B3 alloy as shown in Fig. 4(a). This might be a result that this alloy is presumed to have a higher degree of dense, random packed structure. In contrast, the Mg65Cu25Y7B3 alloy has two small exothermic peaks. In non-isothermal analyses, the dependence of crystallization temperature on heating rate could be used to estimate the associated crystallization activation energy by means of the Kissinger peak shift method [12], Fig. 6. DSC traces showing the melting endotherms obtained from the Mg65Cu25KxY10 Bx (xZ3, 5) alloys at a heating rate of 20 K/min.
Table 1 Thermal properties of the Mg65Cu25-xY10Bx (xZ0, 3, 5) alloys Compositions
Mg65Cu25Y10 Mg65Cu22Y10B3 Mg65Cu20Y10B5
Tg (K)
Tx (K)
DTx (K)
Tmsolid (K)
Tmliquid (K)
Tmliquid
410 410 420
468 476 470
58 66 50
728 715 712
770 738 736
0.532 0.556 0.571
Tg =
g
0.40 0.42 0.41
even smaller melting interval (23 K) indicate that the replacement of Cu by B would result in the apparent narrowing of solid/liquid two phase region and would enhance the GFA of the Mg–Cu–Y amorphous alloys. Also, the single endothermic peak and the narrow melting interval of the Mg65Cu25KxY10Bx with xZ5 demonstrate that this alloy may be very close to the quaternary eutectic composition. The Trg of the optimum Mg65Cu22Y10B3 alloy is 0.556, higher than that for the parent Mg65Cu25Y10 or it corresponds to a 5% increment. The summary of Tg, Tx, DTx, Tm and g (defined as Tx/TgC Tl) [11] is shown in Table 1. From the glass forming g index, the Mg65Cu22Y10B3 alloy again has the highest value of 0.42 than all others, or a 5% increment. Also, as a result of the narrow interval between Tm and Tx, it can contribute to the bulk amorphous formation without a high cooling rate.
lnðb=T 2 Þ ZKEa =RT C constant;
(1)
where b is the heating rate, T is the specific onset temperature, R is gas constant, and Ea is the associated activation energy. Fig. 7 shows the Kissinger’s plots of ln(b/Tx) against 1/Tx, taken from the dependence of crystallization temperature Tx on heating rate b (10, 20, 30 K/min) in the DSC curves. The associated activation energy of the parent Mg65Cu25Y10 amorphous alloy determined from the slope of Kissinger’s plot is 138 kJ/mol, as shown in Fig. 7(a). As a comparison, the activation energy for Mg65Cu22Y10B3 is calculated to be 156 kJ/mol, as compared in Fig. 7(b). The value is about 13% higher than the based alloy. It implies that the substitution of Cu by B would improve the thermal stability and impose a higher energy barrier against crystallization for the Mg-based alloys. The improvement of GFA in the Mg65Cu25KxY10Bx alloys might be a result of multiple factors, including the increment of packing density in supercooled liquid, the larger overall atomic size difference and the higher overall electronegativity difference [13]. The strong affinity for B/Mg and B/Y will suppress the formation of Mg2Cu, MgY and CuY phases, and indirectly enhance the thermal stability of the Mg65Cu25KxY10Bx alloys. Furthermore, both Y and B have a stronger affinity for oxygen than Mg and Cu. A small amount of B addition might be able to suppress the formation of yttrium oxide nuclei, but excessive B would again induce the formation of the B2O3 and YB4. It follows that a critical amount of B addition (w3%) appears to be optimum. A similar effect might also occur in the
Fig. 7. Kissinger plots for the melt spun alloys: (a) Mg65Cu25Y10 (b) Mg65Cu22Y10B3.
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Mg65Cu25Tb10 amorphous alloy [14], where Tb also shows high formation enthalpy to form TbO2. 4. Conclusions The effects of adding interstitial B atoms to the well-known Mg65Cu25Y10 amorphous alloy are examined. The modified Mg65Cu22Y10B3 alloy, with partial replacement of Cu by 3 at% B, exhibits a number of improved properties as compared to Mg65Cu25Y10, for example, a wider supercooled region of 66 K (14% increment versus the original 55 K for Mg65Cu25Y10), a narrower solid/liquid two-phase region of 23 K (43% reduction versus the original 40 K), a higher glass forming g index of 0.42 (5% increment versus the original 0.40), a higher reduced glass transition temperature Trg of 0.556 (5% increment versus the original 0.532), a larger crystallization activation energy of 156 kJ/mol (13% increment versus the original 138 kJ/mol) and an improved hardness of 310 HV (41% increment versus the original 220 HV). This demonastrated that by a proper addition of boron to the parent Mg65Cu25Y10 amorphous alloy, Mg-based amorphous alloys with better thermal stability and mechanical property can be developed.
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