- Email: [email protected]
High specific strength Mg-based bulk metallic glass matrix composite highly ductilized by Ti dispersoid
Materials Science and Engineering A 494 (2008) 299–303
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
High specific strength Mg-based bulk metallic glass matrix composite highly ductilized by Ti dispersoid Makoto Kinaka a , Hidemi Kato b,∗ , Masashi Hasegawa b,1 , Akihisa Inoue b a Department of Materials Science and Engineering, Graduate School of Engineering, Tohoku University, Aramaki aza Aoba 6-6, Sendai 980-8579, Japan b Institute for Materials Research, Tohoku University, Katahira 2-1-1, Sendai 980-8577, Japan
a r t i c l e
i n f o
Article history: Received 9 August 2007 Received in revised form 9 April 2008 Accepted 21 April 2008 Keywords: Mg-based alloy Bulk metallic glass matrix composite High specific strength
a b s t r a c t Mg-based bulk metallic glass matrix composite with hcp-Ti powders was fabricated by casting the mixtures of Mg65 Cu25 Gd10 molten alloy and pure Ti powders into a copper mold. Ti powder was spherical and less than 150 m in diameter, and its volume fraction was controlled from 5 to 40%. Thermal stability of the glassy matrix was maintanied even in the coexistence with the Ti powders. However, Ti dispersoid caused a significant improvement on compressive ductility from 0% plastic deformation for the monolithic glass to 41% plastic deformation for the composite with 40 vol.% Ti powders. This is the first success of synthesizing Mg-based alloys with high ultimate strength of ∼900 MPa level as well as the large plastic deformation of ∼40%, and suggests a novel guideline to develop Mg-based alloys having high specific strength with high ductility. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Recently there has been a strong demand of developing a highstrength and high-ductility material with low specific weight for the maintenance of clean atmosphere on the earth through the saving of energy and other natural resources. Among main engineering metallic alloys, Mg metal has advantages of the lowest specific weight, a large amount of deposits on the earth and its easy reuse. Therefore, many efforts have been done for developing new type materials of Mg-based alloys. Mg-based bulk metallic glasses (BMGs) was found to have good castability as well as high strength about 800 MPa which is about four times higher than AZ91 (Mg90 Al9 Zn1 , a typical high specific strength Mg-based crystalline alloy) [1–11]. However, the Mg-based BMGs do not exhibit appreciable plastic deformation in a uniaxial mode and the improvement on ductility has been strongly requested for the progress in application of Mg-based BMGs. It has been demonstrated that the second phase (i.e., refractory ceramic or metallic particles) dispersed homogeneously in a BMG matrix is effective for improving ductility of BMGs such as Zr-based [12,13] and Cu-based [14,15] alloys under a uniaxial compressive mode. In Mg-based alloys, the first development of Mg-based bulk metallic glass matrix compos-
∗ Corresponding author. Tel.: +81 22 215 2112; fax: +81 22 215 2111. E-mail address: [email protected] (H. Kato). 1 Current address: Graduate School and School of Engineering, Nagoya University, Furo-cho, Chikusa, Nagoya 464-8603, Japan. 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.04.039
ite (BMGMC), a Mg65 Cu7.5 Ni7.5 Zn5 Ag5 Y10 with Fe dispersoid, was reported by Ma et al. in 2003 [16]. Then, many other types have been developed with ceramic and metallic particles, e.g., TiB2 [17], ZrO2 [18], WC [19] and SiC [20], and Nb [21]. However, there has been no data on the formation of Mg-based BMGMC exhibiting high compressive ductility exceeding 10%. Very recently, we have found that the addition of immiscible Ti powders to Mg65 Cu25 Gd10 BMG (specific gravity: 3.7 g/cm3 [4]) causes drastic improvement on plastic deformation up to 41% with maintaining high fracture strength over 850 MPa. This paper intends to the microstructure and excellent mechanical properties of Mg65 Cu25 Gd10 BMGMCs with homogeneously dispersed hcp-Ti powders (specific gravity: 4.5 g/cm3 [22]), and also the mechanism for the remarkable increase in ductility in conjunction with high specific strength by Ti powders.
2. Experimental procedure A Mg65 Cu25 Gd10 (at.%) master alloy was chosen because of its high glass-forming ability leading to a bulk glassy rod of 8 mm in diameter by the conventional copper mold casting technique in Mg65 Cu25 Ln10 (Ln: lanthanide) ternary system. An intermediate ingot was prepared by the induction melting method in an argon atmosphere with pure Mg and Cu–Gd ingot which had been preliminary alloyed by the arc-melting method. Subsequently, crushed pieces of the intermediate Mg65 Cu25 Gd10 ingot and Ti powders with the desired volume fraction was mixed by the induction melting method in a quartz nozzle at a temperature between the
300
M. Kinaka et al. / Materials Science and Engineering A 494 (2008) 299–303
Fig. 1. X-ray diffraction patterns of composite rods of 3 mm in diameter produced by casting the mixtures of Mg65 Cu25 Gd10 molten alloy and the hcp-Ti powders with various volume fractions from 0 to 40% into a copper mold.
liquidus temperature of the pre-alloy (742 K [4]) and melting point of Ti metal (1940 K [22]), followed by casting with a copper mold of inner cavity of 3 mm in diameter in an argon atmosphere. The shape and size of the Ti powders used for the present study were spherical and −100 mesh (less than about 150 m), respectively. The volume fraction of Ti powders to the Mg–Cu–Gd pre-alloy ingot was varied in the range from 0 to 40 vol.%. Structure of the composite alloys was examined by the X-ray diffractometry (XRD) and scanning electron microscopy (SEM) with energy dispersive X-ray (EDX) spectroscopy. Thermal properties associated with the glass transition and crystallization was investigated by a differential scanning calorimeter (DSC) at a heating rate of 0.33 K/s. Mechanical properties were examined under a uniaxial compressive test with an Instron testing machine at the ambient temperature. The dimension of the test specimen was 3 mm in diameter and 6 mm in height (aspect ratio = 2.0), and strain rate was set to be 5.0 × 10−4 s−1 . The fractgraphic examination was carried out by the SEM. 3. Result and discussion Fig. 1 shows X-ray diffraction patterns taken from the transverse cross section of the Mg65 Cu25 Gd10 /Ti composites of 3 mm in diam-
Fig. 2. Scanning electron micrographic image of the transverse cross-section of the cast composite alloy containing 40 vol.% the Ti powders.
Fig. 3. Heat flow curves of the cast Mg65 Cu25 Gd10 bulk metallic glass matrix composites with different volume fractions of Ti powders ranging from 5 to 40% at a heating rate of 0.33 K/s.
eter, together with that of the monolithic Mg65 Cu25 Gd10 metallic glass. The XRD patterns of the composites exhibit sharp peaks diffracted from the hcp-Ti phase in addition to the characteristic broad peaks from the glassy structure. It is worth pointing out that no appreciable diffraction except from the hcp-Ti and glassy phases was observed for all the composites. The Mg65 Cu25 Gd10 alloy has good glass-forming ability which is high enough to suppress precipitation of any crystalline phases even in coexistence with Ti powders. In addition, there is no appreciable change in the position of the main broad peak diffracted from the glassy phase with the volume fraction of Ti powders, indicating that the Mg–Cu–Gd glassy matrix includes little amount of Ti element because of the immiscible behavior between Mg and Ti elements at both the liquid and solid states. In order to clarify the dispersing condition of Ti powders in the glassy matrix, the transverse cross sectional structure of the Mg65 Cu25 Gd10 BMGMC with Ti powders was examined by the SEM. Fig. 2 shows the SEM image of the transverse cross-section of the Mg-based BMGMC with 40 vol.% Ti powders. It is found that the Ti powders maintain its original size and shape (spherical morphology with the diameter less than about 150 m), and disperse very homogeneously in the glassy matrix. This indicates that Ti powders
Fig. 4. True compressive stress–strain curves of the cast Mg65 Cu25 Gd10 bulk metallic glass matrix composites with different volume fractions of Ti powders ranging from 0 to 40% at a strain rate of 5 × 10−4 s−1 .
M. Kinaka et al. / Materials Science and Engineering A 494 (2008) 299–303
301
Fig. 5. (a) Scanning electron micrographic images of fracture surface of the cast composite alloy containing 40 vol.% Ti powders subjected to final fracture after plastic strain of about 40% and (b) energy dispersive X-ray (EDX) spectroscopic mapping of (a).
did not dissolve into the Mg-based glassy matrix. In addition, no distinct segregation of Ti powders is observed. The isolated state of Ti powders surrounded with the glassy phase demonstrates that the Mg-based molten alloy had good castability and high fluidity as well as good wetability with Ti powders. Fig. 3 shows DSC curves of the Mg65 Cu25 Gd10 BMGMCs with different volume fractions of Ti powders. The glass transition temperature (Tg ), crystallization temperature (Tx ) and the exothermic peak temperature for a crystallization reaction (Tp ) are 412, 482 and 488 K, respectively, for the monolithic Mg65 Cu25 Gd10 BMG, and they remain substantially constant over the whole volume fraction range of Ti powders up to 40 vol.%. Therefore, it was confirmed that Ti powders do not cause any appreciable degradation in the thermal stability of the glass and supercooled liquid of the Mg-based matrix alloy, being consistent with the XRD and SEM results. Fig. 4 shows true stress–strain curves of the Mg65 Cu25 Gd10 BMGMCs with different volume fraction of Ti powders under a uniaxial compressive mode at the ambient temperature. The monolithic Mg65 Cu25 Gd10 glassy rod exhibits the Young’s modulus of 49.8 GPa and high fracture strength of 857 MPa in the absence of appreciable plastic deformation. Nearly the same mechanical properties are obtained for the composite with 5 vol.% Ti powders. However, the feature of the mechanical properties changes dras-
tically with further increasing volume fraction of Ti powders. The increase in the volume fraction of the Ti powders from 10 to 40 vol.% causes a significant increase in the compressive plastic deformation (from the elastic limit to the fracture point) from 5% at 10 vol.% to 41% at 40 vol.% through 22% at 20 vol.% and 37% at 30 vol.% as well as a gradual decrease in the yield stress from 800 MPa at 10 vol.% to 470 MPa at 40 vol.%. On the other hand, the Young’ modulus of the composites increases slightly with increasing the volume fraction of Ti powders, and reaches about 62 GPa at 40 vol.%. These are caused by higher ductility, lower yielding stress (103–540 MPa [22]) as well as higher Young’s modulus (∼120 GPa [22]) of Ti powder than those of the Mg-based BMG matrix. It is worth pointing out that the composites with above 10 vol.% Ti powders exhibit distinct work-hardening phenomenon after the yielding because of the work-hardening nature of Ti phase. Therefore, the work hardening rate (d/dε) estimated in a strain range from 0.03 to 0.04 is found to increase with increasing the volume fraction of Ti powder from 1074 MPa at 10 vol.% to 1751 MPa at 40 vol.% through 1258 MPa at 20 vol.% and 1722 MPa at 30 vol.%. The ultimate compressive strength shows a maximum value of 937 MPa at 20 vol.%, followed by nearly the same strength of about 897 MPa at 30 and 40 vol.%, and then about 882 MPa at 10 vol.%. It is particularly noted that the large plastic deformation reaching about 41% at room temperature
Fig. 6. Scanning electron micrographic images of (a) initial test specimen and (b) the outer surface of the cast composite alloy containing 40 vol.% Ti powder subjected to compressive deformation to about 35% strain, and (c) an enlarged image of (b).
302
M. Kinaka et al. / Materials Science and Engineering A 494 (2008) 299–303
Cu–Hf–Ti–Ta [31], Mg-based BMG/Fe [16], Mg-based BMG/TiB2 [17], and Zr-based BMG/ZrC [12,13], which have been reported to exhibit plastic deformation under compression. The Mg65 Cu25 Gd10 BMGMC with 40 vol.% Ti powders is found to have high specific strength of 3.3 × 105 N m/kg with remarkable plastic deformation of 40%. Therefore, the ductile and high specific strength Ti as the filler material is very effective for improving ductility with maintaining the intrinsic high specific strength of the Mg-based BMG. 4. Conclusion
Fig. 7. Specific fracture strength and plasticity of the Mg65 Cu25 Gd10 BMGMC with Ti powders together with typical monolithic BMGs and other types of BMGMCs, i.e., Pt–Cu–Ni–P [23], Zr–Al–Ni–Pd [24,25], Zr–Cu–Al–Pd [26], Zr–Ta–Cu–Ni–Al [27], Zr–Nb–Cu–Ni–Al [28], Zr–Ni–Cu–Ta–Al [29], Ti–Cu–Ni–Sn [30], Cu–Hf–Ti–Ta [31], Mg-based BMG/Fe [16], Mg-based BMG/TiB2 [17], and Zr-based BMG/ZrC [12,13], which have been reported to exhibit plastic deformation.
has not been obtained for monolithic Mg-based BMGs [1–11] as well as Mg-based BMGMCs including any kinds of other crystalline phases [16–21]. In order to clarify the reason for this large compressive plasticity in Mg-based BMGMC, deformation and fracture behaviors of the composite with 40 vol.% Ti powders were examined by the SEM with EDX spectroscopy. Fig. 5(a) and (b) shows the SEM images of fracture surface and the corresponding elemental mapping, respectively. It is observed that the glassy region in the fracture surface consists mainly of a well-developed vein pattern and Ti powders were deformed appreciably, then were cut finally by the shear action at the maximum shear plane. Fig. 6 shows the outer appearance of the initial test specimen (a) and a specimen subjected to 35% plastic deformation (at the lower resolution (b) and higher resolution (c)). In Fig. 6(b) and (c), a large number of shear bands are observed within the glassy region accompanying the significant change in the morphology of Ti powders from the initial sphere to flattened spheroid. These metallographic data reveal that the total plastic deformation of the composite is achieved with slipping of multiple shear-bands in the glassy matrix as well as large plastic deformation in Ti powders. Considering that the yield strength is 103–540 MPa for pure Ti metal [22] and about 850 MPa [4] for Mg65 Cu25 Gd10 metallic glass, the plastic deformation of this composite occurs by the following sequent manner. Dispersed Ti powders, at first, yield so that the stress concentration site is created at the matrix part facing Ti powder and their interface. When the concentrated stress reaches shear-yielding level of the glassy matrix, shear bands begin to arise then to slip into the glassy matrix. Due to the work-hardening nature of Ti powders the composite also exhibited work-hardening behavior. A number of shear bands arose from the interface into the glassy matrix due to coexistence of a number of Ti powders. They are considered to interact, thus, to suppress the fatal instantaneous slip at a single shear band. This allowed a number of shear bands to extend further stably until the final fracture and resulted in the formation of the large plastic deformation in the composite. Fig. 7 summarizes the specific fracture strength and plasticity of the Mg65 Cu25 Gd10 BMGMC with Ti powders together with typical monolithic BMGs and other types of BMGMCs, i.e., Pt–Cu–Ni–P [23], Zr–Al–Ni–Pd [24,25], Zr–Cu–Al–Pd [26], Zr–Ta–Cu–Ni–Al [27], Zr–Nb–Cu–Ni–Al [28], Zr–Ni–Cu–Ta–Al [29], Ti–Cu–Ni–Sn [30],
The Mg65 Cu25 Gd10 bulk metallic glass composite with spherical hcp-Ti powders dispersed homogeneously in the glassy matrix was successfully fabricated by the conventional copper mold casting technique. Diameter of Ti powders was less than 150 m and the volume fraction was controlled from 0 to 40 vol.%. Thermal stability of the glassy matrix did not change with the volume fraction of Ti powders. However, dispersion of Ti powders into the glassy matrix caused a significant increase in the compressive plastic deformation from 0% for the monolithic metallic glass to 41% for the 40 vol.% Ti dispersed composite although the yield strength decreased gradually from 800 MPa at 10 vol.% to 470 MPa at 40 vol.% Ti powders. Ti powders of which yield stress is lower than the glassy matrix induced the stress concentration site at the glassy matrix and facing Ti powders and their interface. These stress concentrated sites created multiple shear bands in the glassy matrix. The interaction among a number of shear bands and the work-hardening nature of Ti powders suppressed the fatal instantaneous slip at a single shear band, thus resulted in the formation of the large plastic deformation as well as high ultimate compressive strength of the composites. This is the first success of synthesizing Mgbased alloys with high ultimate strength of 900 MPa level and the large plastic deformation of 40%, and suggests a novel guideline to develop Mg-based alloys having high specific strength with high ductility. Acknowledgement One of the authors (H.K.) is grateful for the helpful discussion with Prof. H.S. Kim (Chungnam National University, Korea). References [1] A. Inoue, A. Kato, T. Zhang, S.G. Kim, T. Masumoto, Mater. Trans. JIM 32 (1991) 609–616. [2] E.S. Park, H.G. Kang, W.T. Kim, D.H. Kim, J. Non-Cryst. Solids 279 (2001) 154– 160. [3] H. Men, Z.Q. Hu, J. Xu, Scripta Mater. 46 (2002) 699–703. [4] H. Men, D.H. Kim, J. Mater. Res. 18 (2003) 1502–1504. [5] H. Ma, E. Ma, J. Xu, J. Mater. Res. 18 (2003) 2288–2291. [6] E.S. Park, W.T. Kim, D.H. Kim, Mater. Trans. 45 (2004) 2474–2477. [7] H. Ma, Q. Zheng, J. Xu, Y. Li, E. Ma, J. Mater. Res. 20 (2005) 2252–2255. [8] E.S. Park, J.Y. Lee, D.H. Kim, J. Mater. Res. 20 (2005) 2379–2385. [9] G.Y. Yuan, C. Qin, A. Inoue, J. Mater. Res. 20 (2005) 394–400. [10] Q. Zheng, S. Cheng, J.H. Strader, E. Ma, J. Xu, Scripta Mater. 56 (2007) 161–164. [11] E.S. Park, J.S. Kyeong, D.H. Kim, Mater. Sci. Eng. A 449 (2007) 225–229. [12] H. Kato, A. Inoue, Mater. Trans. JIM 38 (1997) 793–800. [13] T. Hirano, H. Kato, Y. Kawamura, A. Matsuo, A. Inoue, Mater. Trans. JIM 41 (2000) 1454–1459. [14] H. Kato, K. Yubuta, D.V. Louzguine, A. Inoue, H.S. Kim, Scripta Mater. 51 (2004) 577–581. [15] H. Kato, D.V. Louzguine, A. Inoue, H.S. Kim, S.I. Hong, J. Metastable Nanocryst. Mater. 15–16 (2003) 161–166. [16] H. Ma, J. Xu, E. Ma, Appl. Phys. Lett. 83 (2003) 2793–2795. [17] Y.K. Xu, H. Ma, J. Xu, E. Ma, Acta Mater. 53 (2005) 1857–1866. [18] J.S.C. Jang, L.J. Chang, J.H. Young, J.C. Huang, C.Y.A. Tsao, Intermetallics 14 (2006) 945–950. [19] P.Y. Lee, C. Lo, J.S.C. Jiang, J.C. Huang, Key Eng. Mater. 313 (2006) 25–29. [20] J. Li, L. Wang, H.F. Zhang, Z.Q. Hu, H. Cai, Mater. Lett. 61 (2007) 2217 –2221. [21] D.G. Pan, H.F. Zhang, A.M. Wang, Z.Q. Hu, Appl. Phys. Lett. 89 (2006) 261904.
M. Kinaka et al. / Materials Science and Engineering A 494 (2008) 299–303 [22] E.A. Brandes, G.B. Brook, Smithells Metals Reference Book, 7th ed., Butterworth–Heinemann Ltd., Oxford, 1992. [23] J. Schroers, W.L. Johnson, Phys. Rev. Lett. 93 (2004) 255506. [24] H. Kato, J. Saida, A. Inoue, Scripta Mater. 51 (2004) 1063–1064. [25] H. Kato, A. Inoue, J. Saida, Appl. Phys. Lett. 85 (2004) 2205–2207. [26] C. Fan, A. Inoue, Appl. Phys. Lett. 77 (2000) 46–48. [27] L.Q. Xing, Y. Li, K.T. Ramesh, J. Li, T.C. Hufnagel, Phys. Rev. B 64 (2001) 180201.
303
¨ ¨ [28] J. Das, W. Loser, U. Kuhn, J. Eckert, S.K. Roy, L. Shultz, Appl. Phys. Lett. 82 (2003) 4690–4692. [29] C. Fan, R.T. Ott, T.C. Hufnagel, Appl. Phys. Lett. 81 (2002) 1020–1022. ¨ [30] G. He, J. Eckert, W. Loser, Acta Mater. 51 (2003) 1621–1631. [31] C.L. Qin, W. Zhang, H. Kimura, A. Inoue, Mater. Trans. JIM 45 (2004) 2936– 2940.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
High specific strength Mg-based bulk metallic glass matrix composite highly ductilized by Ti dispersoid Makoto Kinaka a , Hidemi Kato b,∗ , Masashi Hasegawa b,1 , Akihisa Inoue b a Department of Materials Science and Engineering, Graduate School of Engineering, Tohoku University, Aramaki aza Aoba 6-6, Sendai 980-8579, Japan b Institute for Materials Research, Tohoku University, Katahira 2-1-1, Sendai 980-8577, Japan
a r t i c l e
i n f o
Article history: Received 9 August 2007 Received in revised form 9 April 2008 Accepted 21 April 2008 Keywords: Mg-based alloy Bulk metallic glass matrix composite High specific strength
a b s t r a c t Mg-based bulk metallic glass matrix composite with hcp-Ti powders was fabricated by casting the mixtures of Mg65 Cu25 Gd10 molten alloy and pure Ti powders into a copper mold. Ti powder was spherical and less than 150 m in diameter, and its volume fraction was controlled from 5 to 40%. Thermal stability of the glassy matrix was maintanied even in the coexistence with the Ti powders. However, Ti dispersoid caused a significant improvement on compressive ductility from 0% plastic deformation for the monolithic glass to 41% plastic deformation for the composite with 40 vol.% Ti powders. This is the first success of synthesizing Mg-based alloys with high ultimate strength of ∼900 MPa level as well as the large plastic deformation of ∼40%, and suggests a novel guideline to develop Mg-based alloys having high specific strength with high ductility. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Recently there has been a strong demand of developing a highstrength and high-ductility material with low specific weight for the maintenance of clean atmosphere on the earth through the saving of energy and other natural resources. Among main engineering metallic alloys, Mg metal has advantages of the lowest specific weight, a large amount of deposits on the earth and its easy reuse. Therefore, many efforts have been done for developing new type materials of Mg-based alloys. Mg-based bulk metallic glasses (BMGs) was found to have good castability as well as high strength about 800 MPa which is about four times higher than AZ91 (Mg90 Al9 Zn1 , a typical high specific strength Mg-based crystalline alloy) [1–11]. However, the Mg-based BMGs do not exhibit appreciable plastic deformation in a uniaxial mode and the improvement on ductility has been strongly requested for the progress in application of Mg-based BMGs. It has been demonstrated that the second phase (i.e., refractory ceramic or metallic particles) dispersed homogeneously in a BMG matrix is effective for improving ductility of BMGs such as Zr-based [12,13] and Cu-based [14,15] alloys under a uniaxial compressive mode. In Mg-based alloys, the first development of Mg-based bulk metallic glass matrix compos-
∗ Corresponding author. Tel.: +81 22 215 2112; fax: +81 22 215 2111. E-mail address: [email protected] (H. Kato). 1 Current address: Graduate School and School of Engineering, Nagoya University, Furo-cho, Chikusa, Nagoya 464-8603, Japan. 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.04.039
ite (BMGMC), a Mg65 Cu7.5 Ni7.5 Zn5 Ag5 Y10 with Fe dispersoid, was reported by Ma et al. in 2003 [16]. Then, many other types have been developed with ceramic and metallic particles, e.g., TiB2 [17], ZrO2 [18], WC [19] and SiC [20], and Nb [21]. However, there has been no data on the formation of Mg-based BMGMC exhibiting high compressive ductility exceeding 10%. Very recently, we have found that the addition of immiscible Ti powders to Mg65 Cu25 Gd10 BMG (specific gravity: 3.7 g/cm3 [4]) causes drastic improvement on plastic deformation up to 41% with maintaining high fracture strength over 850 MPa. This paper intends to the microstructure and excellent mechanical properties of Mg65 Cu25 Gd10 BMGMCs with homogeneously dispersed hcp-Ti powders (specific gravity: 4.5 g/cm3 [22]), and also the mechanism for the remarkable increase in ductility in conjunction with high specific strength by Ti powders.
2. Experimental procedure A Mg65 Cu25 Gd10 (at.%) master alloy was chosen because of its high glass-forming ability leading to a bulk glassy rod of 8 mm in diameter by the conventional copper mold casting technique in Mg65 Cu25 Ln10 (Ln: lanthanide) ternary system. An intermediate ingot was prepared by the induction melting method in an argon atmosphere with pure Mg and Cu–Gd ingot which had been preliminary alloyed by the arc-melting method. Subsequently, crushed pieces of the intermediate Mg65 Cu25 Gd10 ingot and Ti powders with the desired volume fraction was mixed by the induction melting method in a quartz nozzle at a temperature between the
300
M. Kinaka et al. / Materials Science and Engineering A 494 (2008) 299–303
Fig. 1. X-ray diffraction patterns of composite rods of 3 mm in diameter produced by casting the mixtures of Mg65 Cu25 Gd10 molten alloy and the hcp-Ti powders with various volume fractions from 0 to 40% into a copper mold.
liquidus temperature of the pre-alloy (742 K [4]) and melting point of Ti metal (1940 K [22]), followed by casting with a copper mold of inner cavity of 3 mm in diameter in an argon atmosphere. The shape and size of the Ti powders used for the present study were spherical and −100 mesh (less than about 150 m), respectively. The volume fraction of Ti powders to the Mg–Cu–Gd pre-alloy ingot was varied in the range from 0 to 40 vol.%. Structure of the composite alloys was examined by the X-ray diffractometry (XRD) and scanning electron microscopy (SEM) with energy dispersive X-ray (EDX) spectroscopy. Thermal properties associated with the glass transition and crystallization was investigated by a differential scanning calorimeter (DSC) at a heating rate of 0.33 K/s. Mechanical properties were examined under a uniaxial compressive test with an Instron testing machine at the ambient temperature. The dimension of the test specimen was 3 mm in diameter and 6 mm in height (aspect ratio = 2.0), and strain rate was set to be 5.0 × 10−4 s−1 . The fractgraphic examination was carried out by the SEM. 3. Result and discussion Fig. 1 shows X-ray diffraction patterns taken from the transverse cross section of the Mg65 Cu25 Gd10 /Ti composites of 3 mm in diam-
Fig. 2. Scanning electron micrographic image of the transverse cross-section of the cast composite alloy containing 40 vol.% the Ti powders.
Fig. 3. Heat flow curves of the cast Mg65 Cu25 Gd10 bulk metallic glass matrix composites with different volume fractions of Ti powders ranging from 5 to 40% at a heating rate of 0.33 K/s.
eter, together with that of the monolithic Mg65 Cu25 Gd10 metallic glass. The XRD patterns of the composites exhibit sharp peaks diffracted from the hcp-Ti phase in addition to the characteristic broad peaks from the glassy structure. It is worth pointing out that no appreciable diffraction except from the hcp-Ti and glassy phases was observed for all the composites. The Mg65 Cu25 Gd10 alloy has good glass-forming ability which is high enough to suppress precipitation of any crystalline phases even in coexistence with Ti powders. In addition, there is no appreciable change in the position of the main broad peak diffracted from the glassy phase with the volume fraction of Ti powders, indicating that the Mg–Cu–Gd glassy matrix includes little amount of Ti element because of the immiscible behavior between Mg and Ti elements at both the liquid and solid states. In order to clarify the dispersing condition of Ti powders in the glassy matrix, the transverse cross sectional structure of the Mg65 Cu25 Gd10 BMGMC with Ti powders was examined by the SEM. Fig. 2 shows the SEM image of the transverse cross-section of the Mg-based BMGMC with 40 vol.% Ti powders. It is found that the Ti powders maintain its original size and shape (spherical morphology with the diameter less than about 150 m), and disperse very homogeneously in the glassy matrix. This indicates that Ti powders
Fig. 4. True compressive stress–strain curves of the cast Mg65 Cu25 Gd10 bulk metallic glass matrix composites with different volume fractions of Ti powders ranging from 0 to 40% at a strain rate of 5 × 10−4 s−1 .
M. Kinaka et al. / Materials Science and Engineering A 494 (2008) 299–303
301
Fig. 5. (a) Scanning electron micrographic images of fracture surface of the cast composite alloy containing 40 vol.% Ti powders subjected to final fracture after plastic strain of about 40% and (b) energy dispersive X-ray (EDX) spectroscopic mapping of (a).
did not dissolve into the Mg-based glassy matrix. In addition, no distinct segregation of Ti powders is observed. The isolated state of Ti powders surrounded with the glassy phase demonstrates that the Mg-based molten alloy had good castability and high fluidity as well as good wetability with Ti powders. Fig. 3 shows DSC curves of the Mg65 Cu25 Gd10 BMGMCs with different volume fractions of Ti powders. The glass transition temperature (Tg ), crystallization temperature (Tx ) and the exothermic peak temperature for a crystallization reaction (Tp ) are 412, 482 and 488 K, respectively, for the monolithic Mg65 Cu25 Gd10 BMG, and they remain substantially constant over the whole volume fraction range of Ti powders up to 40 vol.%. Therefore, it was confirmed that Ti powders do not cause any appreciable degradation in the thermal stability of the glass and supercooled liquid of the Mg-based matrix alloy, being consistent with the XRD and SEM results. Fig. 4 shows true stress–strain curves of the Mg65 Cu25 Gd10 BMGMCs with different volume fraction of Ti powders under a uniaxial compressive mode at the ambient temperature. The monolithic Mg65 Cu25 Gd10 glassy rod exhibits the Young’s modulus of 49.8 GPa and high fracture strength of 857 MPa in the absence of appreciable plastic deformation. Nearly the same mechanical properties are obtained for the composite with 5 vol.% Ti powders. However, the feature of the mechanical properties changes dras-
tically with further increasing volume fraction of Ti powders. The increase in the volume fraction of the Ti powders from 10 to 40 vol.% causes a significant increase in the compressive plastic deformation (from the elastic limit to the fracture point) from 5% at 10 vol.% to 41% at 40 vol.% through 22% at 20 vol.% and 37% at 30 vol.% as well as a gradual decrease in the yield stress from 800 MPa at 10 vol.% to 470 MPa at 40 vol.%. On the other hand, the Young’ modulus of the composites increases slightly with increasing the volume fraction of Ti powders, and reaches about 62 GPa at 40 vol.%. These are caused by higher ductility, lower yielding stress (103–540 MPa [22]) as well as higher Young’s modulus (∼120 GPa [22]) of Ti powder than those of the Mg-based BMG matrix. It is worth pointing out that the composites with above 10 vol.% Ti powders exhibit distinct work-hardening phenomenon after the yielding because of the work-hardening nature of Ti phase. Therefore, the work hardening rate (d/dε) estimated in a strain range from 0.03 to 0.04 is found to increase with increasing the volume fraction of Ti powder from 1074 MPa at 10 vol.% to 1751 MPa at 40 vol.% through 1258 MPa at 20 vol.% and 1722 MPa at 30 vol.%. The ultimate compressive strength shows a maximum value of 937 MPa at 20 vol.%, followed by nearly the same strength of about 897 MPa at 30 and 40 vol.%, and then about 882 MPa at 10 vol.%. It is particularly noted that the large plastic deformation reaching about 41% at room temperature
Fig. 6. Scanning electron micrographic images of (a) initial test specimen and (b) the outer surface of the cast composite alloy containing 40 vol.% Ti powder subjected to compressive deformation to about 35% strain, and (c) an enlarged image of (b).
302
M. Kinaka et al. / Materials Science and Engineering A 494 (2008) 299–303
Cu–Hf–Ti–Ta [31], Mg-based BMG/Fe [16], Mg-based BMG/TiB2 [17], and Zr-based BMG/ZrC [12,13], which have been reported to exhibit plastic deformation under compression. The Mg65 Cu25 Gd10 BMGMC with 40 vol.% Ti powders is found to have high specific strength of 3.3 × 105 N m/kg with remarkable plastic deformation of 40%. Therefore, the ductile and high specific strength Ti as the filler material is very effective for improving ductility with maintaining the intrinsic high specific strength of the Mg-based BMG. 4. Conclusion
Fig. 7. Specific fracture strength and plasticity of the Mg65 Cu25 Gd10 BMGMC with Ti powders together with typical monolithic BMGs and other types of BMGMCs, i.e., Pt–Cu–Ni–P [23], Zr–Al–Ni–Pd [24,25], Zr–Cu–Al–Pd [26], Zr–Ta–Cu–Ni–Al [27], Zr–Nb–Cu–Ni–Al [28], Zr–Ni–Cu–Ta–Al [29], Ti–Cu–Ni–Sn [30], Cu–Hf–Ti–Ta [31], Mg-based BMG/Fe [16], Mg-based BMG/TiB2 [17], and Zr-based BMG/ZrC [12,13], which have been reported to exhibit plastic deformation.
has not been obtained for monolithic Mg-based BMGs [1–11] as well as Mg-based BMGMCs including any kinds of other crystalline phases [16–21]. In order to clarify the reason for this large compressive plasticity in Mg-based BMGMC, deformation and fracture behaviors of the composite with 40 vol.% Ti powders were examined by the SEM with EDX spectroscopy. Fig. 5(a) and (b) shows the SEM images of fracture surface and the corresponding elemental mapping, respectively. It is observed that the glassy region in the fracture surface consists mainly of a well-developed vein pattern and Ti powders were deformed appreciably, then were cut finally by the shear action at the maximum shear plane. Fig. 6 shows the outer appearance of the initial test specimen (a) and a specimen subjected to 35% plastic deformation (at the lower resolution (b) and higher resolution (c)). In Fig. 6(b) and (c), a large number of shear bands are observed within the glassy region accompanying the significant change in the morphology of Ti powders from the initial sphere to flattened spheroid. These metallographic data reveal that the total plastic deformation of the composite is achieved with slipping of multiple shear-bands in the glassy matrix as well as large plastic deformation in Ti powders. Considering that the yield strength is 103–540 MPa for pure Ti metal [22] and about 850 MPa [4] for Mg65 Cu25 Gd10 metallic glass, the plastic deformation of this composite occurs by the following sequent manner. Dispersed Ti powders, at first, yield so that the stress concentration site is created at the matrix part facing Ti powder and their interface. When the concentrated stress reaches shear-yielding level of the glassy matrix, shear bands begin to arise then to slip into the glassy matrix. Due to the work-hardening nature of Ti powders the composite also exhibited work-hardening behavior. A number of shear bands arose from the interface into the glassy matrix due to coexistence of a number of Ti powders. They are considered to interact, thus, to suppress the fatal instantaneous slip at a single shear band. This allowed a number of shear bands to extend further stably until the final fracture and resulted in the formation of the large plastic deformation in the composite. Fig. 7 summarizes the specific fracture strength and plasticity of the Mg65 Cu25 Gd10 BMGMC with Ti powders together with typical monolithic BMGs and other types of BMGMCs, i.e., Pt–Cu–Ni–P [23], Zr–Al–Ni–Pd [24,25], Zr–Cu–Al–Pd [26], Zr–Ta–Cu–Ni–Al [27], Zr–Nb–Cu–Ni–Al [28], Zr–Ni–Cu–Ta–Al [29], Ti–Cu–Ni–Sn [30],
The Mg65 Cu25 Gd10 bulk metallic glass composite with spherical hcp-Ti powders dispersed homogeneously in the glassy matrix was successfully fabricated by the conventional copper mold casting technique. Diameter of Ti powders was less than 150 m and the volume fraction was controlled from 0 to 40 vol.%. Thermal stability of the glassy matrix did not change with the volume fraction of Ti powders. However, dispersion of Ti powders into the glassy matrix caused a significant increase in the compressive plastic deformation from 0% for the monolithic metallic glass to 41% for the 40 vol.% Ti dispersed composite although the yield strength decreased gradually from 800 MPa at 10 vol.% to 470 MPa at 40 vol.% Ti powders. Ti powders of which yield stress is lower than the glassy matrix induced the stress concentration site at the glassy matrix and facing Ti powders and their interface. These stress concentrated sites created multiple shear bands in the glassy matrix. The interaction among a number of shear bands and the work-hardening nature of Ti powders suppressed the fatal instantaneous slip at a single shear band, thus resulted in the formation of the large plastic deformation as well as high ultimate compressive strength of the composites. This is the first success of synthesizing Mgbased alloys with high ultimate strength of 900 MPa level and the large plastic deformation of 40%, and suggests a novel guideline to develop Mg-based alloys having high specific strength with high ductility. Acknowledgement One of the authors (H.K.) is grateful for the helpful discussion with Prof. H.S. Kim (Chungnam National University, Korea). References [1] A. Inoue, A. Kato, T. Zhang, S.G. Kim, T. Masumoto, Mater. Trans. JIM 32 (1991) 609–616. [2] E.S. Park, H.G. Kang, W.T. Kim, D.H. Kim, J. Non-Cryst. Solids 279 (2001) 154– 160. [3] H. Men, Z.Q. Hu, J. Xu, Scripta Mater. 46 (2002) 699–703. [4] H. Men, D.H. Kim, J. Mater. Res. 18 (2003) 1502–1504. [5] H. Ma, E. Ma, J. Xu, J. Mater. Res. 18 (2003) 2288–2291. [6] E.S. Park, W.T. Kim, D.H. Kim, Mater. Trans. 45 (2004) 2474–2477. [7] H. Ma, Q. Zheng, J. Xu, Y. Li, E. Ma, J. Mater. Res. 20 (2005) 2252–2255. [8] E.S. Park, J.Y. Lee, D.H. Kim, J. Mater. Res. 20 (2005) 2379–2385. [9] G.Y. Yuan, C. Qin, A. Inoue, J. Mater. Res. 20 (2005) 394–400. [10] Q. Zheng, S. Cheng, J.H. Strader, E. Ma, J. Xu, Scripta Mater. 56 (2007) 161–164. [11] E.S. Park, J.S. Kyeong, D.H. Kim, Mater. Sci. Eng. A 449 (2007) 225–229. [12] H. Kato, A. Inoue, Mater. Trans. JIM 38 (1997) 793–800. [13] T. Hirano, H. Kato, Y. Kawamura, A. Matsuo, A. Inoue, Mater. Trans. JIM 41 (2000) 1454–1459. [14] H. Kato, K. Yubuta, D.V. Louzguine, A. Inoue, H.S. Kim, Scripta Mater. 51 (2004) 577–581. [15] H. Kato, D.V. Louzguine, A. Inoue, H.S. Kim, S.I. Hong, J. Metastable Nanocryst. Mater. 15–16 (2003) 161–166. [16] H. Ma, J. Xu, E. Ma, Appl. Phys. Lett. 83 (2003) 2793–2795. [17] Y.K. Xu, H. Ma, J. Xu, E. Ma, Acta Mater. 53 (2005) 1857–1866. [18] J.S.C. Jang, L.J. Chang, J.H. Young, J.C. Huang, C.Y.A. Tsao, Intermetallics 14 (2006) 945–950. [19] P.Y. Lee, C. Lo, J.S.C. Jiang, J.C. Huang, Key Eng. Mater. 313 (2006) 25–29. [20] J. Li, L. Wang, H.F. Zhang, Z.Q. Hu, H. Cai, Mater. Lett. 61 (2007) 2217 –2221. [21] D.G. Pan, H.F. Zhang, A.M. Wang, Z.Q. Hu, Appl. Phys. Lett. 89 (2006) 261904.
M. Kinaka et al. / Materials Science and Engineering A 494 (2008) 299–303 [22] E.A. Brandes, G.B. Brook, Smithells Metals Reference Book, 7th ed., Butterworth–Heinemann Ltd., Oxford, 1992. [23] J. Schroers, W.L. Johnson, Phys. Rev. Lett. 93 (2004) 255506. [24] H. Kato, J. Saida, A. Inoue, Scripta Mater. 51 (2004) 1063–1064. [25] H. Kato, A. Inoue, J. Saida, Appl. Phys. Lett. 85 (2004) 2205–2207. [26] C. Fan, A. Inoue, Appl. Phys. Lett. 77 (2000) 46–48. [27] L.Q. Xing, Y. Li, K.T. Ramesh, J. Li, T.C. Hufnagel, Phys. Rev. B 64 (2001) 180201.
303
¨ ¨ [28] J. Das, W. Loser, U. Kuhn, J. Eckert, S.K. Roy, L. Shultz, Appl. Phys. Lett. 82 (2003) 4690–4692. [29] C. Fan, R.T. Ott, T.C. Hufnagel, Appl. Phys. Lett. 81 (2002) 1020–1022. ¨ [30] G. He, J. Eckert, W. Loser, Acta Mater. 51 (2003) 1621–1631. [31] C.L. Qin, W. Zhang, H. Kimura, A. Inoue, Mater. Trans. JIM 45 (2004) 2936– 2940.