Crystallisation behaviour of Cu60Zr30Ti10 bulk glassy alloy

Materials Science and Engineering A 375–377 (2004) 744–748

Crystallisation behaviour of Cu60 Zr30 Ti10 bulk glassy alloy Masayuki Kasai a,∗ , Eiichiro Matsubara b , Junji Saida a , Masao Nakayama a , Katsuhito Uematsu c , Tabo Zhang b , Akihisa Inoue b a

Inoue Superliquid Glass Project, Exploratory Research for Advanced Technology, Japan Science and Technology Corporation, Yagiyamaminami 2-1-1, Sendai 982-0807, Japan b Institute for Materials Research (IMR), Tohoku University, Sendai 980-8577, Japan c Graduate School, Tohoku University, Sendai 980-8577, Japan

Abstract The crystallisation behaviour in the Cu60 Zr30 Ti10 alloy was studied. The as-quenched Cu60 Zr30 Ti10 alloy consists of homogeneously dispersed nanocrystals in the glassy matrix. Two metastable crystalline phases appear prior to the precipitation of equilibrium phases, which correspond to the three exothermic peaks in the DSC curve. After the first exothermic reaction, the nanometer-sized CuZr phase is precipitated in the glassy matrix. In the second exothermic peak, the CuZr phase transforms to a metastable cubic phase with a large lattice parameter of 1.265 nm. Finally, this metastable phase is decomposed to Cu10 Zr7 , Cu8 Zr3 and Cu3 Ti2 equilibrium phases. The Cu10 Zr7 and Cu8 Zr3 phases contain Ti while the Cu3 Ti2 phase contains Zr. © 2003 Published by Elsevier B.V. Keywords: Metallic glasses; X-ray diffraction (XRD); Phase transformation; Cu-based alloy

1. Introduction Recently, new Cu-based alloys exhibiting good mechanical strength of 2000–2160 MPa have been found in Cu–(Hf or Zr)–Ti systems [1]. The alloys containing more than 50 at.% Cu are important for further development as engineering materials. These alloys have a unique feature in that they do not satisfy the three empirical rules for high glass-forming ability [2]. The authors have previously reported microstructures of the melt-spun Cu60 (Zr or Hf)30 Ti10 alloys and have also suggested that the nanocrystalline particles exist in the glassy matrix even in the as-quenched state [3]. In these studies, it has been reported that the Cu-rich nanocrystalline cubic phases are directly formed during the melt-quenching process in these alloys and the high stability of the Cu-rich nanocrystalline cubic phases leads to a coexistence with the glassy phase on a nanometer scale. In this study, we will examine the thermal characteristics and crystallisation behaviour of the Cu60 Zr30 Ti10 glassy

∗ Corresponding author. Tel.: +81-22-243-7661; fax: +81-22-243-7616. E-mail address: [email protected] (M. Kasai).

0921-5093/$ – see front matter © 2003 Published by Elsevier B.V. doi:10.1016/j.msea.2003.10.092

alloys with special attention to the structure of metastable crystalline phases.

2. Experimental Cu60 Zr40−x Tix (x = 0 and 10 at.%) alloys were prepared by arc-melting the mixtures of pure Cu (99.9 mass%), Zr (99.9 mass%) and Ti (99.9 mass%) metals in a purified argon atmosphere, followed by melt-spinning. Thermal characteristics of these alloys were examined by differential scanning calorimetry (DSC) at a heating rate of 0.67 K/s in a purified (99.999 wt.%) argon atmosphere. The as-quenched Cu60 Zr30 Ti10 alloy was annealed for 30–600 s in the temperature range of 753–1000 K in an infrared furnace under vacuum. The heating rate was chosen similar to DSC measurements at 0.67 K/s and the residual pressure of the furnace was kept to be less than 3 × 10−5 Pa during annealing. The structure of the annealed samples was examined by X-ray diffraction (XRD) with Cu K␣ radiation. Lattice parameters and volume fractions of the equilibrium phases were determined by the Rietveld method [4]. The microstructure in as-quenched state and primary metastable phase was examined by field emission-transmission electron microscopy (FE-TEM) with an accelerating voltage of 300 kV (JEOL

M. Kasai et al. / Materials Science and Engineering A 375–377 (2004) 744–748

JEM-3000F). The compositional analysis was performed by nanobeam energy dispersive X-ray spectroscopy (EDS).

3. Results and discussion Fig. 1 shows the DSC curves of the melt-spun Cu60 Zr40−x Tix (x = 0 and 10 at.%) alloys. The glass transition temperature (Tg ) is at 719 K for Cu60 Zr40 and at 712 K for Cu60 Zr30 Ti10 . Three exothermic peaks are observed in the ternary alloy. The onset temperature of the first-stage crystallisation (Tx ) corresponding to the lower exothermic peak is 760 K for the binary alloy and 749 K for the ternary alloys. The Tg and Tx decrease by addition of Ti and the temperature interval between the first and second exothermic peaks is 55 K for the ternary alloy. Fig. 2 shows the bright-field TEM image (a), selected-area electron diffraction pattern (SADP) taken from the region of 1 ␮m in diameter (b) and high-resolution TEM image (c) of the as-quenched Cu60 Zr30 Ti10 alloy. In the bright-field TEM image, the diffuse contrast is seen over the whole region. It is seen in the high-resolution TEM image that the origin of this contrast is attributed to a nanocrystalline phase of about 3 nm in diameter, which is embedded in the glassy matrix. The lattice constants of the nanocrystalline phase which precipitated in the Cu60 Zr30 Ti10 alloy has been reported to be a = 0.45 nm [4]. In order to determine the structure of the primary phase which precipitated from the as-quenched state, we examined the structure of the Cu60 Zr30 Ti10 alloy annealed for 600 s at 753 K corresponding to the transformation due to the first exothermic peak. Fig. 3(a) shows the XRD patterns of the Cu60 Zr30 Ti10 alloy in the as-quenched and annealed states. Only diffused peaks are observed in the annealed sample in Fig. 3(a). This fact indicates that a mixture of residual glassy and nanocrystalline phases is formed in the first exothermic reaction. The nanometer-sized crystalline phase is identified from the differential X-ray intensity profile between inten-

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sities of the annealed and as-quenched samples in Fig. 3(b). These diffraction peaks can be indexed as the CsCl-type bcc ¯ [5]. CuZr phase (Pm3m) We also examined the first metastable phase in the Cu60 Zr30 Ti10 alloy by TEM. The bright-field TEM image (a), SADP (b) and NBD pattern (c) are shown in Fig. 4. Very fine particles of diameters less than 10 nm are observed over the whole area and embedded in the residual glassy matrix. The precipitates have a nearly spherical morphology and are homogeneously distributed. The SADP exhibits broad ring pattern and many fine diffraction spots. In order to analyze the structure of the particles, we examined nanobeam diffraction patterns of the particles with a beam diameter of about 2 nm. Although a couple of diffraction patterns are overlapped in Fig. 4(c), a clear sixfold symmetry is observed. The lattice constant calculated from the diffraction spots is 0.328 nm, which is equal to that of the bcc CuZr structure. The TEM observation result agrees well with that obtained by the XRD measurement. We investigated the transformation from the first metastable bcc CuZr phase to the second metastable phase in the Cu60 Zr30 Ti10 alloy. Fig. 5 shows the XRD patterns of the Cu60 Zr30 Ti10 alloy annealed for 30–300 s at 803 K corresponding to the phase transformation due to the second exothermic peak. As seen in Fig. 5, some diffuse peaks appear in the sample annealed for 60 s, and the crystalline phase progressively grows in the sample annealed for 300 s. All these peaks are indexed by assuming the symmetry of the simple-cubic structure with the lattice constant of 1.265 nm. It appears that the primary nanometer-sized bcc CuZr phase transforms into the cubic phase with the large lattice parameter of 1.265 nm at the second exothermic peak. Fig. 6 shows the XRD pattern of the Cu60 Zr30 Ti10 alloy annealed for 600 s at 1000 K. As it is expected from the ternary Cu–Zr–Ti phase diagram [6], the metastable large cubic phase is transformed into Cu10 Zr7 , Cu8 Zr3 and Cu3 Ti2 phases at the final stage of crystallisation. Thus it is concluded that the transformation process of the Cu60 Zr30 Ti10 alloy is expressed as follows: glassy phase + cubic phase → bcc CuZr + residual glassy phase

Exothermic

→ large cubic phase → Cu10 Zr 7 + Cu8 Zr 3 + Cu3 Ti2 . Tg

Cu60Zr40 Tx Cu60Zr30Ti10

Tg

Tx 700

800

900

Temperature (K) Fig. 1. DSC curves of the melt-spun Cu60 Zr40−x Tix (x = 0 and 10 at.%) alloys.

We also applied the Rietveld method [4] in order to refine the lattice parameters and volume fractions of the equilibrium phases. The results are summarised in Table 1. The lattice parameters of Cu10 Zr7 are shrunk by 1.1% for a, 0.90% for b and 1.4% for c, while those of Cu8 Zr3 are slightly shrunk by 0.029% for a, 0.12% for b and 0.044% for c. On the other hand, the lattice parameters of Cu3 Ti2 are slightly shrunk by 0.09% for a and extended by 0.16% for c. The change of the lattice parameters in these crystals is quantitatively explained by the substitution of the Zr sites by the different amounts of Ti atoms in each crystal. The volume fractions obtained in the present analysis are

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Fig. 2. Bright-field TEM image (a), selected-area electron diffraction pattern (SADP) (b) and high-resolution TEM image (c) of the Cu60 Zr30 Ti10 glassy alloy. Table 1 Crystallographic systems, the refined lattice parameters in the present work and the volume fractions of each equilibrium phase in the Cu60 Zr30 Ti10 alloy annealed at 1000 K for 600 s Crystal

Crystallographic system (space group no.)

Refined lattice parameters (nm)

Lattice [7] parameters (nm)

Percentage of change (%)

Volume fraction (%)

Cu10 Zr7

Orthorhombic (Aba2 no.41)

a = 9.2430 b = 9.2324 c = 12.495

a = 9.3466 b = 9.3163 c = 12.6729

−1.100 −0.900 −1.404

82.7

a = 7.8715 b = 8.1536 c = 9.9804

a = 7.8693 b = 8.1547 c = 9.9848

−0.028 −0.012 −0.044

14.0

a = 3.1272 c = 13.973

a = 3.130 c = 13.95

−0.089 0.160

3.3

Cu8 Zr3

Cu3 Ti2

Orthorhombic (Pnma no.62)

Tetragonal (P4/nmm no.129)

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Fig. 3. XRD patterns of the Cu60 Zr30 Ti10 alloy and that of the sample annealed for 600 s at 753 K (a) and the differential X-ray intensity,  intensity (b).

consistent with those evaluated from the ternary phase diagram [6]. The lattice parameter (1.265 nm) of the second metastable cubic phase is nearly equal to that (1.267 nm) of c-axis in the Cu10 Zr7 main equilibrium phase, while the lattice parameters of a- and b-axes are about 3/4 as those of the cubic phase. It implies that the atomic structure of the large cubic phase resembles that of the main equilibrium

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Fig. 5. XRD patterns of the Cu60 Zr30 Ti10 alloy annealed at 803 K for 300 s (a), 60 s (b) and 30 s (c).

phase Cu10 Zr7 . By assuming the structural similarity of these phases, the heat released of the third exothermic peak in the ternary alloy is expected to be smaller than the other two peaks. It has been reported that a number of bulk glassy alloys satisfy the three empirical rules for high glass-forming

Fig. 4. Bright-field TEM image (a), selected-area electron diffraction pattern (SADP) (b) and nanobeam electron diffraction (NBD) pattern (c) of the Cu60 Zr30 Ti10 alloy annealed for 600 s at 753 K.

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4. Conclusions

Fig. 6. XRD pattern of the equilibrium phases of the Cu60 Zr30 Ti10 alloy annealed for 600 s at 1000 K.

ability, i.e. (1) multicomponent system, (2) large atomic size mismatches above 12% and (3) suitable negative heats of mixing [2]. The Cu–Zr–Ti systems do not completely satisfy the above criteria. That is, Ti does not have distinct chemical affinity with Zr. This implies that a novel structure in the as-quenched state is different from those in other bulk-forming glassy alloys [8,9]. It was shown by the microanalyses using nanobeam diffraction and nanobeam EDS that the as-quenched Cu60 Zr30 Ti10 alloy is in coexistence with nanocrystalline particles with a Cu-rich cubic structure [3]. The average compositions in the nanocrystalline and glassy phases are Cu68.8 Zr25.7 Ti5.5 and Cu56.1 Zr35.6 Ti8.3 , respectively. It was elucidated that the Zr and Ti atoms are enriched and the Cu element is rejected in the glassy matrix. We also suggest that the stability of Cu-rich nanocrystalline cubic phase is attributed to the Ti element. Ti has an atomic size between Cu and Zr and does not have distinct chemical affinity with Zr. This fact seems to cause the increase in packing density of the glassy phase and the suppression of the rearrangement of the constitutional elements, which are the main factors for the thermal stability of nanocrystalline phases in the as-quenched state. The primary crystallisation phase corresponding to the first exothermic peak in the DSC curve is a bcc CuZr phase in the Cu60 Zr30 Ti10 alloy. The composition change during the precipitation of the first metastable nanometer-sized CuZr phase in the Cu60 Zr30 Ti10 ternary alloy was examined by nanobeam EDS. The beam diameter for the nanobeam EDS was approximately 2 nm. An average composition at five data points in the CuZr phase is Cu59.7 Zr32.7 Ti7.6 , which is almost equal to Cu57.3 Zr34.7 Ti8.1 in the residual glassy phase. The similarity of the compositions of the glassy and CuZr phases suggests that the nucleation of the CuZr phase takes place in the glassy phase, and the residual glassy phase crystallises into the bcc CuZr phase in the coexistence with the Cu-rich nanocrystalline cubic phase.

We investigated crystallisation behaviour of the Cu60 Zr40−x Tix (x = 0 and 10 at.%) glassy alloys. Two metastable crystalline phases appear prior to precipitation of the equilibrium phases in the ternary alloy. After the first exothermic reaction, the Cu60 Zr30 Ti10 alloy is transformed to a mixture of a residual glassy phase and nanometer-sized metastable CuZr crystalline phase. During the second metastable phase in the Cu60 Zr30 Ti10 alloy, the residual glassy phase transforms to a cubic phase with a large lattice parameter of 1.265 nm. Finally, this metastable phase decomposes to Cu10 Zr7 , Cu8 Zr3 and Cu3 Ti2 equilibrium phases. In comparison with the lattice parameters of the equilibrium crystalline phases in the ternary systems, it is suggested that the Cu10 Zr7 and Cu8 Zr3 phases contain Ti and that the Cu3 Ti2 phase contains Zr. The Cu60 Zr30 Ti10 alloy, even in the as-quenched state, evolves the homogeneous dispersion of nanocrystals in the glassy matrix, which may be the origin for the much higher tensile strength compared with that for the corresponding glassy single phase alloy. Nanocrystalline materials have been attracting much attention as advanced materials. There have been a number of reports on the formation of nanocrystalline phases embedded in the glassy matrix by appropriate heat treatment of the as-quenched glassy alloys [10,11]. In the present alloy, however, the nanocrystals and glass composite materials are formed even in the as-quenched state. The use of this advantage, which cannot be obtained for the unstable glassy alloys, could provide a new and simple way to control the microstructure of nanocrystalline phase embedded in the glassy matrix.

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