Intermetallics 19 (2011) 1474e1478
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Mechanical behaviour of Zr-La-Cu-Ni-Al glass-based composites Shantanu V. Madge a, *, Parmanand Sharma b, Dmitri V. Louzguine-Luzgin a, A. Lindsay Greer a, c, Akihisa Inoue a, b a
World Premier International Research Center (WPI-AIMR), Tohoku University, 2-1-1 Katahira, Aoba-Ku, Sendai 980-8577, Japan Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Aoba-Ku, Sendai 980-8577, Japan c Department of Materials Science & Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ, UK b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 10 November 2010 Received in revised form 16 May 2011 Accepted 19 May 2011 Available online 14 June 2011
A range of phase-separating Zr-La-Cu-Ni-Al alloys has been cast. The Zr-based alloys show a dispersion of La-based, partially amorphous spheres in a Zr-based glassy matrix while the La-based alloys consist of Zr-based, crystalline spheres in a La-based glassy matrix. Compression testing reveals embrittlement in the former alloys, but toughening in the latter. The observed behaviour is explained by considering the toughness of the reinforcing phase, in comparison to the glassy matrix. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: B. Glasses, metallic mechanical properties at ambient temperature B. Mechanical properties at ambient temperature
1. Introduction Bulk metallic glasses (BMGs) possess attractive properties like high strength and a low modulus. However, monolithic BMGs have limited plasticity because failure occurs via highly localised deformation inside a single shear band [1]. The key to improving the toughness of BMGs lies in avoiding the propagation of a single shear band to failure, and in ensuring that deformation occurs via the formation of multiple shear bands, which increases the macroscopic plastic strain. This is often achieved by dispersing crystals in a glassy matrix that block shear band propagation and lead to their proliferation [1]. Phase separation in bulk metallic glasses (BMGs) is an issue that has seen much attention in the recent past, partly in order to explain the nanocrystallisation observed in some compositions [2e4]. Amorphous phase separation could also be used to obtain glass-glass or glass-crystal composites with various microstructural length-scales and possibly higher toughness. A few studies have shown that the presence of two amorphous phases does hinder the propagation of shear bands in a variety of metallic glasses and thereby enhances the toughness [5e10]. Embrittlement has also been observed in some systems, though, e.g. in Cu-(Zr,Hf)(Gd,Y)-Al glasses [11,12].
* Corresponding author. E-mail address:
[email protected] (S.V. Madge). 0966-9795/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2011.05.015
Like Y, La is immiscible with Zr [13], and Kündig et al. [4] could synthesise a phase-separating glass, La27.5Zr27.5Al25Cu10Ni10, in the form of melt-spun ribbons. In the present work, the Zr-La-Cu-Ni-Al system is explored further. Both, Zr-based and La-based phaseseparating alloys have been cast in bulk form. However, glass-glass composites have not been obtained. Instead, the Zr-based alloys contain a dispersion of partially crystalline La-based spheres in a Zr-based glassy matrix and vice versa. Although the original aim of getting glass-glass composites was not realised, these glasscrystal composites have a wide range of microstructural length scales (100 nme5 mm) and thus, their mechanical properties are of some interest. Embrittlement has been noted in the Zr-rich series, whereas some toughening is seen in the La-rich alloys and the behaviour is explained in terms of the deformation mechanism. An interesting aspect of the Zr-based alloys is that, upon fracture, the dispersed La-based spheres appear to undergo melting. This finding is supported by a quantitative estimate of the heating associated with fracture of the Zr-based alloys. 2. Experimental procedures Ingots of the well known glass-formers Zr55Cu30Ni5Al10 and La55Al25Cu10Ni10 were first prepared by arc-melting a mixture of the pure elements (>99.9 wt.% pure). Pieces of these ingots were mixed in different proportions (10% and 20% by weight) and arc-melted together to arrive at the following compositions:Zr50.6La4.4Ni5.4Cu28.4Al11.2 (A);
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Fig. 1. SEM back-scattered electron images of (a) Zr46La9Ni5.8Cu26.7Al12.5 (alloy B) and (b) La41.8Zr13.3Ni8.8Cu14.8Al21.3 (alloy D). Alloy B contains La-based spheres in a Zr-based glassy matrix and vice versa for alloy D. (c) XRD patterns from the present series of alloys. The Zr-based alloys are X-ray amorphous, whereas peaks corresponding to a mixture of Zr-rich intermetallics (mostly Zr3Al and AlNiZr) are detectable in the La-based alloys C and D.
Zr46La9Ni5.8Cu26.7Al12.5(B); La48.2Zr6.8Ni9.4Cu12.5Al23.1 (C) and La41.8Zr13.3Ni8.8Cu14.8Al21.3(D). All the compositions are given in atomic %. The alloys were induction-melted and cast into Cu moulds to get 2 mm diameter rods. The amorphicity of the samples was checked using CuKa X-ray radiation. Microstructural studies were done using a Hitachi S-4800 field-emission gun scanning electron microscope (FEGSEM) and compositional analysis was done using electron probe micro-analysis (EPMA) in a JEOL Superprobe JXA-8621 MX instrument. Transmission electron microscopy (TEM) observations were made using a JEOL 2010 instrument with a LaB6 filament, operated at 200 kV. TEM specimens were prepared by the usual procedure of dimpling followed by cold-stage ion-milling in a Fischione ion-miller operated at 3 kV. Specimens for compression testing were cut from the 2 mm rods and were carefully polished to have a final length of 4 mm, using a special fixture designed to ensure that their ends were parallel. The testing was carried out at an engineering strain rate of 5 104 s1. Scratch-testing on polished specimens was done to study the interaction between shear bands and any second phase. This was done using a movable stage and a Vickers diamond pyramid indenter at a sliding speed of approximately 1 mm/s with an applied load of 2 kg.
3. Results and discussion 3.1. Structural characterisation Fig. 1a,b shows the typical phase-separated microstructures seen in the as-cast specimens. The Zr-based alloys (A and B) consist of La-based spheres dispersed in a Zr-based matrix and vice versa for alloys C and D. Table 1 shows the compositions of the phases, and their approximate volume fractions, deduced from the micrographs. Consistent with [4], spheres of all sizes, from 100 nm to about 5 mm are seen and the phase separation also involves Ni and Cu, i.e. areas rich in Zr are also rich in Ni and areas rich in La are rich in Cu. From the X-ray diffraction patterns (Fig. 1c), the Zr-based alloys appear to be amorphous. Fig. 2aec shows the TEM brightfield images obtained from alloy B. The Zr-based matrix does not show any crystals and its corresponding selected-area electron diffraction pattern (SAD) shows a typical amorphous halo. However, the La-based spheres are largely crystalline (a mixture of LaeCu intermetallics), although some of them are partially amorphous, evident from the diffuse ring seen in their SAD patterns (Fig. 2c). In case of the La-based alloys, the Zr-based spheres that
Table 1 Compositions of the alloys and the two phases formed upon phase separation. Alloy
Matrix
Spheres
Monolithic Zr55Cu30Ni5Al10 Monolithic La55Al25Cu10Ni10 A:Zr50.6La4.4Ni5.4Cu28.4Al11.2 B:Zr46La9Ni5.8Cu26.7Al12.5 C:La48.2Zr6.8Ni9.4Cu12.5Al23.1 D:La41.8Zr13.3Ni8.8Cu14.8Al21.3
e e Zr57La0.3Ni5Cu25.6Al12.1(92.5 vol.%) Zr55.6La0.5Ni6Cu24.4Al13.5 (85 vol. %) La53.8Zr3.0Ni7.4Cu13.3Al22.5 (92.5 vol%) La53Zr3.6Ni5.8Cu16.8Al20.8 (85 vol. %)
e e La47.2Zr2.4Ni1.1Cu39.7Al9.6 (7.5 vol.%) La48.1Zr2.8Ni0.9Cu40Al8.2 (15 vol.%) Zr48.4La2Ni19.2Cu4.9Al25.5(7.5 vol.%) Zr49.8La1.8Ni18.5Cu6.7Al23.3 (15 vol.%)
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Fig. 2. (a) TEM bright field image from as-cast Zr46La9Ni5.8Cu26.7Al12.5 (alloy B), showing a dispersion of La-based spheres in a Zr-based glassy matrix. (b) Selectedarea electron diffraction (SAD) pattern from the Zr-based matrix shows a typical amorphous halo. (c) SAD pattern from the La-based sphere shows crystalline reflections in addition to the glassy halo.
form are completely crystalline, amply evident from the crystalline peaks in the XRD traces (Fig. 1c) that arise from a mixture of hexagonal (hP9) AlNiZr and cubic Zr3Al (cP4). It is worth noting that the La (or Zr)-based spheres that form in the present alloys do not undergo further phase separation to form ‘nested structures’, unlike the Zr-La-Cu-Ni-Al alloy studied by Kündig et al. [4], where each of the initially-formed phases separates further into Zr- and La-rich regions. What is crucial is the starting composition of the alloy. The alloy studied in [4] has equal amounts of Zr and La, i.e. La27.5Zr27.5Cu10Ni10Al25 and consequently, the alloy must be deep inside the miscibility gap. In contrast, alloys in the present study (Table 1) are either toward the Zr- or the Larich end, and consequently, there would be little driving force for further phase-separation within the spheres that have already precipitated from the matrix. Fig. 3 shows the DSC data for the present series of alloys. For all alloys, the matrix shows a clear glass transition and crystallisation, whereas no peaks corresponding to the crystallisation of the dispersed spheres are visible, confirming the XRD and TEM data.
formation. Apparently, the embrittlement in Cu46Zr37Y10Al7 is akin to relaxation-induced embrittlement, seen, for example, in FeeNieB glasses [16] where the material does not form shear bands and failure proceeds through cleavage instead. This is distinct from the present situation, where shear bands do form and the question must be what leads to a change in fracture mechanism from mode II to mode I. Previous work on oxygen-induced embrittlement in a Cubased BMG [17] shows that such a change in the fracture mode is because of flaws in a high-strength material, e.g. oxygen-containing brittle dendrites in a Cu49Hf42Al9 glass, through which a propagating shear band can pass easily. La-based glasses are soft and brittle compared to Zr-based BMGs [14]. Hence, it is likely that the La-based spheres in Fig. 1a might crack upon loading and essentially act like flaws in the Zr-based glassy matrix, accounting for the observed embrittlement. Scratch-testing offers a way of observing the interaction between shear bands and any second-phase particles in a glass [18]. As shown in Fig. 5d, propagating shear bands (around a scratch made using a diamond indenter) in the Zr-based matrix indeed pass through the La-based particles, unobstructed. An intriguing feature about the fracture surface of alloy B (Fig. 5c) is the unexpected cavities seen between the white, La-based particles and the Zr-based glassy matrix. A strong possibility is that the La-based spheres melt during mode I crack propagation in the Zr-based matrix, and the cavities form on their re-solidification. The temperature rise in the La-based spheres upon cracking of the
a
6
exo.
5 4
Heat Flow (W/g)
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Alloy B
3
Alloy A
2 1 0
Zr55Cu30Ni5Al10
-1
3.2. Mechanical properties
100
150
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250
300
350
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450
500
550
600
o
Temperature ( C)
b
8
exo. 6
Alloy D Heat Flow (W/g)
3.2.1. Zr-based alloys Fig. 4a shows the compressive stress-strain curves for Zr55Cu30Ni5Al10 and the Zr-based alloys A and B. The monolithic glass shows an elastic strain-to-failure of w2% and a fracture stress of 1860 25 MPa. Alloys A and B fail at a lower stress (1600 and 1320 MPa respectively) and the scatter is significantly higher, upto 100 MPa. Fig. 5a shows the fracture surface of an alloy A specimen, as an example. Two regions, X and Y, can be distinguished, whose relative proportions vary depending on the specimen. Region X is at 45 to the axis and shows shear band vein patterns (Fig. 5b), and region Y shows nano-scale dimples (Fig. 5c). As is now well known [14] the length scale of the fractographic features reflects the toughness of a BMG, with micron-scale vein patterns being typical of tough glasses and sub-micron scale dimples being characteristic of brittle glasses. A tough BMG can also show nano-scale features if the fracture mode is not shear (mode II) but local tensile failure, i.e. mode I [15]. Thus, Fig. 5a shows that failure in alloy A is by mode II to begin with, but quickly shifts to mode I, resulting in brittle fracture. Embrittlement on phase separation was also reported by Park et al. [11] in a Cu46Zr37Y10Al7 BMG, a major difference being that this alloy shows brittle fracture, without any evidence for shear band
4
2
Alloy C
0
-2
La55Al25Cu10Ni10
-4 100
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o
Temperature ( C) Fig. 3. DSC traces from (a) the Zr-based alloys and (b) La-based alloys.
500
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Since the La-based spheres in Fig. 5c are about same size as the plastic zone ahead of a mode I crack in the Zr-based matrix (0.25 mm), it may be safely assumed that the La-based spheres will be uniformly heated through their cross-section, i.e. thermal diffusion through the spheres is not expected to play a major role. The temperature rise in a La-based sphere can be estimated using
G$A ¼ m$Cp $DT
(4)
where A is the cross-sectional area of the sphere and m its mass. The density of the La-based spheres is taken to be 6 g/cm3, based on the literature values for similar compositions [22]. For the 0.5 mm diameter La-based spheres in Fig. 5c, DT is expected to be 4743 K, sufficiently above the melting temperature of w723 K (see Fig. 3b). Thus, the cavities around the La-based spheres in Fig. 5c can be accounted for by the melting of those spheres during fracture of the glassy matrix.
Fig. 4. Compressive stress-strain data for the present series of (a) Zr-based alloys and (b) La-based alloys. Compared to Zr55Cu30Ni5Al10, alloys A and B fail at a lower, and irreproducible stress. In contrast, strengthening (and toughening) is achieved in the Labased alloys, C & D.
Zr-based glass can be estimated from the fracture energies, in the following way: The mode I fracture toughness of Zr-based BMGs ranges from 20 to 80 MPa m1/2 depending on the experimental conditions [19]. Assuming KIc w20 MPa m1/2 and Young’s modulus, E ¼ 96 GPa [20], the fracture energy for the matrix,
Gmatrix ¼ K 2 =E ¼ 4:16 kJ=m2
(1)
The fracture energy of a La55Al25Cu10Ni5Co5 bulk glass is reported to be 0.5 kJ/m2 [19]. This may be used as an estimate of the fracture energy for the present La-based spheres (Gspheres), although in reality, Gspheres is likely to be lower since the spheres are largely crystalline. Assuming that mechanical energy is converted into heat, the net energy, G, available for heating the Labased spheres can be written as
G ¼ Gmatrix Gspheres ¼ 3:7 kJ=m2
(2)
The specific heat of crystalline La55Al25Cu10Ni10 is given by Lu et al. [21] as
Cp ¼ 3R þ 5:09 103 T þ 1:624 105 T 2
(3)
For a temperature range of 300 Ke723 K (the melting point), Cp varies from 27.9 to 37.1 J/mol.K. For the present calculation, Cp is assumed to be 37.1 J/mol.K (0.39 J/g.K).
Fig. 5. A set of secondary electron images for the present Zr-La-Cu-Ni-Al alloys. (a) The fracture surface of an alloy A specimen shows two regions, X and Y. (b) A closer view of region X showing micron-scale, shear band vein patterns. (c) Region Y shows nano-scale dimples and the white particles are the same La-based spheres seen in Fig. 1a. The circled region highlights the cavities between spheres and the Zr-based glassy matrix. (d) Results of scratch-testing on alloy B. The dotted circle shows that shear bands in the glassy matrix readily pass through the La-based spheres. (e) Alloy D fails gracefully, by shear along a 45 plane and breaks into two pieces only. (f) Multiple shear offsets are seen on the sides of a fractured specimen of alloy D. (g) Scratch-testing shows that the Zr-based, crystalline spheres in alloy D serve as obstacles to shear bands in the La-based glass. The dotted circles show examples of sites where the shear bands are forced to travel around the Zr-based spheres.
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3.2.2. La-based alloys Unlike the Zr-based alloys, the behaviour of the La-based compositions is more encouraging. Fig. 4b shows the compressive stress-strain curves for the La-based series of alloys. The monolithic La55Al25Cu10Ni10 glass fails at an average stress of 750 MPa, but with considerable scatter (50 MPa) within specimens. Alloys C and D, on the other hand, show a higher strength, significantly less scatter (20 MPa) and although the plastic strain is still low, fractography reveals toughening in these alloys. The La55Al25Cu10Ni10 glass fails in the expected brittle manner by bursting into many pieces (not shown). In contrast, alloys C and D fail by shear at w45 to the axis (Fig. 5e) and break into two pieces only. Importantly, the sides of the specimen show multiple shear offsets (Fig. 5f), indicating an improved toughness. Scratch-testing reveals that the Zr-based spheres indeed serve as obstacles to shear band propagation (Fig. 5g). Available data show that Zr-based glasses are about twice as hard as La-based glasses [14] and crystallisation can be expected to increase the hardness. Hence it is likely that a dispersion of the harder Zr-based particles in the La-based glassy matrix can deflect shear-bands and cause their proliferation. Although some toughening is seen in alloys C&D, the plastic strain-to-failure is still limited, unlike, say Zr-based BMGs reinforced with WC particles [23], which show 3e7% strain. The key difference between these two composites is the toughness of the matrix itself e the Labased glassy matrix is brittle, unlike the Zr-based matrix in [23]. In case of La-based glassy composites, large plasticity was noted in composites employing a ductile (Ti) reinforcement [24]. Thus, a grand challenge in this area would be designing an alloy that forms a homogeneous glass on quenching, and phase-separates on annealing (so that the length-scale and the volume fractions of the two phases can be independently controlled) and then one of the phases forms ductile crystals on continued annealing. 4. Conclusions In summary, a range of phase-separating Zr-La-Cu-Ni-Al alloys has been synthesised. The Zr-based compositions contain La-based, partially amorphous, spherical particles dispersed in a Zr-based glassy matrix. The composite microstructure shows a lower toughness than the monolithic Zr-based glass, because the brittle
La-based spheres serve as flaws in the otherwise tough glassy matrix. The La-based spheres are seen to melt upon fracture of the composite. An analysis based on the fracture energies corroborates this observation, as the estimated temperature rise is about 4700 K. The La-based series of alloys show a dispersion of substantially crystalline Zr-based spheres in a La-based glassy matrix and compression testing reveals an improved toughness compared to the monolithic La-based glass. This is attributable to the higher strength of the Zr-based spheres, which aids in blocking shear band propagation to an extent, although the achieved plastic strain is still limited. References [1] Schuh CA, Hufnagel TC, Ramamurty U. Acta Mater 2007;55:4067e109. [2] Madge SV, Alexander DTL, Greer AL. J Non-Cryst Solids 2003;317:23e9. [3] Liu W, Johnson WL, Schneider S, Geyer U, Thiyagarajan P. Phys Rev B 1999;59: 11755e9. [4] Kündig AA, Ohnuma M, Ping DH, Ohkubo T, Hono K. Acta Mater 2004;52: 2441e8. [5] Du XH, Huang JC, Chen HM, Chou HS, Lai YH, Hsieh KC, et al. Intermetallics 2009;17:607e13. [6] Concustell A, Mattern N, Wendrock H, Kuehn U, Gebert A, Eckert J, et al. Scripta Mater 2007;56:85e8. [7] Oh JC, Ohkubo T, Kim YC, Fleury E, Hono K. Scripta Mater 2005;53:165e9. [8] Van Steenberge N, Concustell A, Sort J, Das J, Mattern N, Gebert A, et al. Mater Sci Eng A 2008;491:124e30. [9] Lee SC, Huh MY, Kim HJ, Lee JC. Mater Sci Eng.A 2008;485:61e5. [10] Ren YL, Zhu RL, Sun J, You JH, Qiu KQ. J Alloys Compd 2010;493:L42e46. [11] Park ES, Kim DH. Acta Mater 2006;54:2597e604. [12] Park ES, Kyeong JS, Kim DH. Scripta Mater 2007;57:49e52. [13] Massalski TB. Binary alloy phase diagrams, vol. 3. ASM International; 1990. 2444. [14] Ki XK, Zhao DQ, Pan MX, Wang WH, Wu Y, Lewandowski JJ. Phys Rev Lett 2005;94:125510 (1e4). [15] Jiang MQ, Ling Z, Meng JX, Dai LH. Phil Mag 2008;88:407e26. [16] Ocelík V, Diko P, Csach K, Hajko V, Bengus VZ, Tabachnikova ED, et al. Mater Sci 1987;22:3732e6. [17] Madge SV, Wada T, Louzguine-Luzgin DV, Greer AL, Inoue A. Scripta Mater 2009;61:540e3. [18] Matthews DTA, Ocelík V, de Hosson J. Th M Mat Sci Eng 2007;A471:155e64. [19] Xu J, Ramamurty U, Ma E. JOM 2010;62:10e8. [20] Ashby MF, Greer AL. Scripta Mater 2006;54:321e6. [21] Lu ZP, Hu X, Li Y. Intermetallics 2000;8:477e80. [22] Dong TY, Yang B, He JP, Zhang Y. Acta Metall Sin 2009;45:232e6. [23] Choi-Yim H, Busch R, Köster U, Johnson WL. Acta Mater 1999;47:2455e62. [24] Madge SV, Sharma P, Louzguine-Luzgin DV, Greer AL, Inoue A. Scripta Mater 2010;62:210e3.
