Thermal stability and crystallization kinetics of Cu-reinforced Cu47Ti33Zr11Ni8Si1 metallic glass composite powders synthesized by ball milling: the effect of particulate reinforcement

Intermetallics 13 (2005) 833–840 www.elsevier.com/locate/intermet

Thermal stability and crystallization kinetics of Cu-reinforced Cu47Ti33Zr11Ni8Si1 metallic glass composite powders synthesized by ball milling: the effect of particulate reinforcement S. Venkataramana,*, E. Rozhkovac, J. Eckerta,b, L. Schultza, D.J. Sordeletc,d a IFW Dresden, Institut fu¨r Metallische Werkstoffe, Postfach 27 00 16, D-01171 Dresden, Germany Technische Universita¨t Darmstadt, FB 11 Material- und Geowissenschaften, FG Physikalische Metallkunde, Petersenstrasse 23, D-64287 Darmstadt, Germany c Material and Engineering Physics Program, Ames Laboratory (USDOE), Iowa State University, Ames, IA 50014, USA d Department of Materials Science and Engineering, Iowa State University, Ames, IA 50014, USA b

Received 7 October 2004; received in revised form 13 January 2005; accepted 13 January 2005 Available online 19 March 2005

Abstract The thermal stability and crystallization kinetics of Cu47Ti33Zr11Ni8Si1 gas atomized powder (GAP) and composite powder reinforced with 25 vol.% nanosized Cu second phase particles introduced by ball milling was studied by differential scanning calorimetry in the mode of continuous heating and isothermal annealing. Ball milling leads to successful incorporation of Cu particles within the GAP matrix and hence offers itself as a useful technique to prepare composite microstructures. In the case of continuous heating, both the glass transition temperature Tg and the crystallization temperature Tx display a strong dependence on the heating rate. The activation energy, as determined by the Kissinger equation, for the first order transformation was found to be almost identical for the GAP (3.53G0.12 eV) and composite powders (3.59G0.06 eV). The isothermal transformation kinetics were modeled by the Johnson–Mehl–Avrami (JMA) equation. The values of the JMA exponent imply that the crystallization of the metallic glass GAP as well as the composite powder is governed by diffusioncontrolled three-dimensional growth. The activation energy of crystallization derived from isothermal annealing experiments is nearly the same for the GAP (3.69G0.20 eV) and the composite powders (3.78G0.10 eV). The addition of Cu particles does not noticeably affect the crystallization kinetics of the amorphous matrix composites. q 2005 Elsevier Ltd. All rights reserved. Keywords: B. Glasses, metallic; C. Mechanical alloying and milling; F. Calorimetry; F. Diffraction (electron, neutron and X-ray); F. Microscopy, various

1. Introduction Metallic glasses are promising candidates for use as matrices in metal-matrix-composites due to their high strength and high elastic strain limit [1,2]. Second phase metallic or ceramic particles homogenously dispersed in a glassy amorphous matrix hinder the propagation of shear bands and initiate multiple shearing events [3–5]. This not only improves the toughness of the material [2] but also * Corresponding author. Tel.: C49 351 4659 687; fax: C49 351 4659 541. E-mail address: [email protected] (S. Venkataraman).

0966-9795/$ - see front matter q 2005 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2005.01.010

enhances the ductility of the metallic glass matrix composites [4]. The improved ductility opens up the possibility to overcome the so far limited applications of the monolithic bulk metallic glasses [6] that fail catastrophically in a brittle manner [7–10]. The production of metallic glass matrix composites by liquid state techniques imposes restrictions on the particle size and the volume fraction of the reinforcements added [3,11]. The glass-forming ability is also affected by the addition of second phase particles in the melt due to dissolution and chemical reactions occurring in the melt at the liquid/solid interface [7]. Alternatively, ball milling/ mechanical alloying allows preparation of glassy composites thereby circumventing the above-mentioned difficulties [12–15]. Glass formation is achieved through a solid

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state reaction governed by intermixing and mechanically activated diffusion processes [13,15,16]. Consolidation of powders into bulk specimens is possible due to the characteristic viscosity minimum in the supercooled liquid region [5,10,14–19]. Recent studies [14,15] on brass reinforced bulk metallic glass composites synthesized by warm extrusion indicate that the addition of brass induces macroscopic plasticity after yielding with values up to 2%. Upon yielding, the ductile reinforcement seems to plastically deform first and the load is then transferred to the surrounding glass matrix. This initiates shear bands at the reinforcement/matrix interface. However, the propagation of shear bands is found to be restricted between the reinforcement fibres. This leads to more homogeneous distribution of plasticity resulting to higher strain to failure and complete avoidance of catastrophic failure. However before commencing on a mechanical property study, it is necessary to first establish if the stability of amorphous GAP matrices are compromised by introducing ductile metallic second phases on a fine length scale. Cu47Ti33Zr11Ni8Si1 is one of the typical bulk glassforming alloys, which exhibits excellent glass-forming ability and a wide supercooled liquid region before crystallization. The addition of Si to the Cu–Ti–Ni–Zr alloy improves the glass-forming ability and enlarges the width of the supercooled liquid region before crystallization [11]. Moreover, an enhanced plasticity compared to its Si-free counterpart has been reported for this alloy produced by copper mold casting [20]. Crystallization studies have been carried out on ball milled metallic glass matrix composites containing carbide particles as a second phase [21]. However, no work has been reported on the crystallization kinetics of Cu 47Ti33Zr 11Ni8 Si 1 glass reinforced with metallic particles. Since the amorphous state is essentially a metastable one, it inherently possesses the possibility of transforming into a more stable crystalline state. The promising properties of metallic glasses, i.e. high hardness and strength, have been found to deteriorate on crystallization [22]. Understanding the micro mechanisms of crystallization is, therefore, a prerequisite for most applications as the stability against crystallization determines their potential use as novel structural materials. Also, a detailed analysis of the thermal stability and crystallization kinetics of milled glassy and composite powders is the key to explore the suitability of ball milling as an effective tool for processing of bulk metallic glass composites. In this paper, we report on the formation of glassy Cu47Ti33Zr11Ni8Si1 composite powders reinforced with 25 vol.% Cu particles by ball milling. In particular, the effect of the second phase particles on the non-isothermal and isothermal crystallization kinetics is investigated by differential scanning calorimetry. The local Avrami exponent was also determined to explain the details of the nucleation and growth behavior in the crystallization process.

2. Experiment Spherical Cu47Ti33Zr11Ni8Si1 gas-atomized powder (GAP) was prepared by high-pressure gas atomization. The details of the powder synthesis have been reported elsewhere [23]. The nominal size of the GAP is 45 mm. Ball milling of the GAP blended with 25 vol.% Cu (99.9%, 10 mm size) was done in a Retsch 4000 planetary mill at a rotational velocity of 100 rpm using hardened steel milling tools and a ball-to-powder ratio of 15:1 for 20 h. All powders were handled in a glove box under purified argon atmosphere (!1 ppm of O2 and H2O). Inductively coupled plasma optical emission spectroscopy (ICP-OES) revealed deviations of about G0.6 at.% from the nominal composition and iron impurities of less than 0.5 at.% due to wear debris from the milling tools. The amount of oxygen in the milled samples was determined by hot extraction (C436 LECO analyzer) to be about 1.7 at.%. Structural characterization was performed by X-ray diffraction (XRD) using a Philips PW1050 diffractometer (Co Ka radiation). Microscopic investigations were carried out using a scanning electron microscopy (JEOL 6400) equipped with an in situ energy dispersive X-ray spectrometer (EDX). Further microstructural characterization was done using a field emission gun scanning electron microscope (LEO GEMINI 1531). The crystallization kinetics of the amorphous alloy and the composite powder was characterized by continuous heating and isothermal annealing in a differential scanning calorimeter (PerkinElmer DSC-7) under flowing high purity Ar. In the case of continuous heating a set of differential scanning calorimetry (DSC) plots was recorded at heating rates ranging from 10 to 80 K/min. For isothermal analysis, the samples were first heated to a fixed temperature (between 718 K and 733 K) at a heating rate of 40 K/min and held for a certain period of time until completion of crystallization. Al pans were used for the continuous heating and isothermal annealing experiments. The DSC system was calibrated for temperature and enthalpy by using zinc and indium standards. For each individual sample, two successive DSC runs were recorded successively. The second run of the specimen serves as a baseline. Subtraction of this baseline from the first run realized the correction for the apparatus specific baseline shift. The liquidus temperature was determined using a Netzsch DSC 404 operated at a heating rate of 20 K/min.

3. Results and discussion 3.1. X-ray diffraction Fig. 1 shows the XRD patterns of the gas atomized powder and the composite powder after 20 h of ball milling. The XRD pattern for the gas-atomized powder (Fig. 1(b)) displays an amorphous pattern. There are no crystalline reflections from the starting metallic elements and the diffraction pattern is typical of an amorphous material [24].

S. Venkataraman et al. / Intermetallics 13 (2005) 833–840

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Cu

Co K α

Intensity (a.u.)

(c) (b)

(a) 20

40

60

80

100

Scattering angle, 2θ Fig. 1. X-ray diffraction patterns of (a) pure copper, (b) Cu47Ti33Zr11Ni8Si1 GAP and (c) composite powder containing 25 vol.% Cu and milled for 20 h.

Fig. 3. Scanning electron micrograph for composite showing the typical layered structure of a ball milled material.

The XRD pattern of powder blended with 25 vol.% Cu and milled for 20 h (Fig. 1(c)) shows the presence of crystalline reflections of the principal Cu peaks superimposed on the broad diffuse scattering maxima of the amorphous phase. Fig. 1(a) shows the reflections of pure annealed Cu for reference. There is no change in the position of the scattering vector of the amorphous maximum. The positions of the Cu peaks match those of pure Cu and the lattice parameter is calculated as 0.3616 nm thereby indicating a shift of 0.004 nm after 20 h of ball milling, the lattice parameter of pure unmilled Cu being 0.3612 nm. This is a very minute change in the lattice parameter as evidenced by the essentially identical XRD patterns between the starting annealed Cu and the composite powders. The grain size of the Cu particles is estimated to be 43 nm after 20 h of ball milling as estimated from the peak broadening of the X-ray peaks by the Williamson–Hall method [25].

are spherical. Crystalline inclusions are also seen in some powder particles. The origin of these inclusions lies in the pre-alloying step where elemental constituents are heated and melted in a graphite crucible prior to casting [23]. Fig. 3 shows the microstructural features of the composite powder containing 25 vol.% Cu milled for 20 h. The powder particles display a layered structure that is formed as a result of the ball milling and the incorporation of the second phase Cu particles. The matrix appears uniform and the Cu particles are quite homogeneously distributed in the matrix. The Cu particles are no longer spherical but are elongated and flattened due to the severe deformation during the ball– powder collision events upon milling resulting in flaky and plate-like structures.

3.2. Microstructure Fig. 2 shows a scanning electron image of the Cu47Ti33 Zr11Ni8Si1 powder. It is obvious that most of the particles

Fig. 2. Scanning electron micrograph of Cu47Ti33Zr11Ni8Si1 GAP.

3.3. Crystallization behavior on continuous heating Fig. 4(a) shows the continuous heating DSC plots recorded at different heating rates for Cu47Ti33Zr11Ni8Si1 gas atomized powder while Fig. 4(b) shows the corresponding scans for the composite powder. All DSC traces exhibit an endothermic event characteristic of a glass transition (the glass transition temperature, Tg, was defined as the onset of the endothermic DSC event) and an extended supercooled liquid region before the onset of crystallization (the crystallization temperature, Tx, was defined as the onset temperature of the first exothermic event). For example, taking the DSC scan for the Cu47Ti33Zr11Ni8Si1 GAP recorded at 20 K/min reveals that the onset temperature of the glass transition, Tg, is 691 K and the onset of crystallization, Tx, is 752 K. Hence, the width of the supercooled liquid region, defined as, DTxZTxKTg, is 61 K for this heating rate. This value is larger than the value of 46 K reported for Cu47Ti34Zr11Ni8 metallic glass produced by copper mold casting [26] thereby indicating that the addition of Si increases the extent of the supercooled liquid region of this alloy. However, the width of the supercooled liquid region is slightly smaller compared to that of as cast

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(a)

20 K/min Exo Down

10 K/min 20 K/min

Heat Flow (a.u.)

Heat flow (a.u.)

30 K/min 40 K/min 60 K/min

80 K/min

TSolidus

TLiquidus

Endo Up

500

550

600

650

700

750

800

850

Temperature, K

1100

1150

1200

Temperature (K)

(b) 10 K/min

Fig. 5. High temperature DSC scan for Cu47Ti33Zr11Ni8Si1 GAP.

20 K/min

influence the crystallization temperature and, hence, increases the width of the supercooled liquid region. As expected, for all the DSC measurements, Tg and Tx and shift to higher temperatures with increasing heating rate, i.e. both the crystallization and the glass transition display a dependence on the heating rate during continuous heating. The activation energy for crystallization Ec can be evaluated by means of the Kissinger equation [30]

40 K/min 60 K/min

80 K/min

lnðF=Tp2 Þ Z KE=RTp C constant

Endo Up

500

550

600

650

700

750

800

850

Temperature, K Fig. 4. Constant heating rate DSC scans for (a) Cu47Ti33Zr11Ni8Si1 GAP, (b) composite powder.

bulk samples of same composition, reported earlier (66 K) [11] but is higher than the one reported recently for the same alloy (57 K) [27]. The comparisons are made with DSC data at 20 K/min with Ref. [11] and at 40 K/min with Ref. [27]. The difference in the extent of supercooled liquid region could be ascribed due to the slight compositional differences of the respective samples and the presence of crystalline inclusions in the gas atomized powder. The reduced glass transition temperature, Trg, is a value to estimate the glass forming ability of alloys and represents the ratio of the glass transition temperature and liquidus temperature. The liquidus temperature for the gas-atomized powder was determined using a high temperature DSC and is shown in Fig. 5. The reduced glass transition temperature is found out to be 0.60. This value is also similar to the value reported earlier [11] and indicates a high glass formability [28,29]. For the composite powder, the onset temperature of glass transition is 688 K and the onset of crystallization is 752 K at a heating rate of 20 K/min, yielding a width of the supercooled liquid region of 64 K. Thus, the addition of Cu decreases the glass transition temperature, does not

(1)

where F is the heating rate and R the gas constant; Tp stands for the crystallization peak temperature. By plotting lnðF=Tp2 Þ vs. 1/Tp, an approximately straight line is obtained, as shown in Fig. 6. The correlation coefficient is better than 0.999. From the slope of the straight line the activation energy for crystallization is obtained. The activation energy values calculated for the gas-atomized amorphous powder (3.53G0.12 eV) and the composite powder (3.59G0.06 eV) are much higher than that of Zr55Cu30Al10Ni5 (2.38 eV) [31] –8.5 GAP COMPOSITE

–9.0

ln(φ/Tp2) (K–1min–1)

Heat flow (a.u.)

30 K/min

–9.5

–10.0

–10.5 EGAP –11.0

: 3.53 ± 0.12 eV

ECOMPOSITE : 3.59 ± 0.06 eV 1.28

1.29

1.30

1.31

1000/Tp

1.32

1.33

1.34

(K–1)

Fig. 6. Kissinger plots for Cu47Ti33Zr11Ni8Si1 GAP and composite.

S. Venkataraman et al. / Intermetallics 13 (2005) 833–840

(a)

Endothermic heat flow (a.u.)

and that of Zr41Ti14Cu12.5Ni10Be22.5 (2 eV) [32] indicating that the Cu-based metallic glass as well as the composite are more resistant to crystallization compared to Zr-based alloys. The values obtained are in good agreement with those obtained for the Cu47Ti33Zr11Ni8Si1 bulk metallic glass obtained by copper mold casting (3.69 eV) [11]. The marginal differences between these values could be attributed to slightly a different composition of the glassy matrix and due to the presence of crystalline inclusions in the gas atomized powders. However, there is no significant difference between the activation energies of the GAP and the composite. Apparently, the interfaces between the GAP and the Cu particles do not act as heterogeneous nucleation sites, which generally decrease the activation barrier for crystallization [11]. This is of significant importance in the synthesis of bulk materials by subsequent consolidation processes such as hot pressing and warm extrusion.

837

718 K 720 K 723 K 725 K

730 K

728 K

733 K

0

10

20

30

Time (min) (b)

The isothermal DSC investigations of the GAP and the composite powder containing Cu particles were done at different temperatures between 718 and 733 K and the corresponding DSC plots are shown in Fig. 7(a) and (b), respectively. All DSC traces exhibit a single exothermic peak after passing a certain incubation period. This corresponds to the first crystallization event in the constant heating rate scans. The incubation time as well as the time required for the complete transformation become shorter for annealing at higher temperatures. This can be ascribed to the higher atomic mobility at higher temperatures, which causes concentration fluctuations necessary for large-scale crystallization to set in. It is assumed that the transformed volume fraction x up to any time t is proportional to the fractional areas of the exothermic peak. The crystallized volume fraction during the isothermal process can be accurately determined by measuring the area of the exothermic peak. The crystallized fraction, xc(t, T), after time t at a constant temperature T was derived from the curves in Fig. 7(a) and (b) by assuming that xc is proportional to the integrated enthalpy, ðt ðN xc ðt; TÞ Z hðt; TÞdt= hðt; TÞdt: (2) 0

0

The kinetics of phase transition is usually analyzed in terms of the generalized theory of phase transition [33] xc ðt; TÞ Z 1 K expðKkðt K tÞn Þ;

(3)

where k is a reaction rate constant, t is the incubation time and n is the Avrami exponent whose value can vary from 1 to 4 [34]. Eq. (3) can be rewritten as the well-known Johnson– Mehl–Avrami (JMA) Equation [35]: ln½ln 1=ð1 K xÞ Z n ln k C n lnðt K tÞ:

(4)

Fig. 8(a) and (b) show the plots of ln[ln 1/(1Kx)] and n ln(tKt) for the gas-atomized starting powder and the ball-milled composite. For 0.10%xc%0.85, the data are

Endothermic heat flow (a.u.)

3.4. Crystallization behavior during isothermal annealing 718 K

720 K 723 K 725 K 728 K 730 K

733 K

0

10

20

30

Time (min) Fig. 7. Isothermal DSC traces for (a) Cu47Ti33Zr11Ni8Si1 GAP, (b) composite powder.

almost on a straight line with a correlation coefficient better than 0.998. The Avrami exponent n and the reaction rate constant k can be calculated from the slopes and intercepts of the lines. The temperature dependence of the Avrami exponent n and the incubation time t are summarized in Tables 1 and 2. The values of n range from 2.8 to 3.5 for the GAP while it ranges from 2.7 to 3.4 for the composite powder. The value of the Avrami exponent relates to the operating crystallization mechanism [33]. A more detailed interpretation [36] of the crystallization kinetics suggests that the Avrami exponent n can be expressed as n Z a C bc;

(5)

where a is the nucleation index (aZ0 for zero nucleation rate, 0!a!1 for decreasing nucleation rate with time, aZ1 for constant nucleation rate and aO1 for increasing nucleation rate) [37], b is the dimensionality of the growth (with values of 1, 2 or 3), and c, the growth index (cZ1 for interfacecontrolled growth and cZ0.5 for diffusion-controlled growth). Crystallization occurs by long-range diffusion of

733 K

730 K

728 K

725 K

718 K

(a) 1.0

723 K

S. Venkataraman et al. / Intermetallics 13 (2005) 833–840

720 K

838

0.5

Table 2 Avrami exponent and incubation time at different temperatures for composite powder Temperature (K)

Incubation time (min)

Avrami exponent (0.1%Xc%0.85)

718 720 723 725 728 730 733

6.10 5.30 2.90 1.83 1.46 1.06 0.76

3.09 3.32 2.74 2.96 3.40 3.05 3.23

ln(-ln(1-Xc))

0.0 –0.5 –1.0 –1.5 –2.0 –2.5 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.2 2.4 2.6 2.8 3.0

718 K

0.5

ln(-ln(1-Xc))

0.0 –0.5 –1.0 –1.5 –2.0 –2.5 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.2 2.4 2.6 2.8 3.0

Nloc Z vln ½Kln ð1 K xÞ=vln ðt K tÞ:

ln(t-τ) (min) Fig. 8. Avrami plot for (a) Cu47Ti33Zr11Ni8Si1 GAP, (b) composite powder.

atoms in the supercooled liquid and, hence, is diffusioncontrolled. So using a value of cZ0.5 in Eq. (5) allows to derive the values of b at different annealing temperatures yielding 3. The Avrami exponent n gives detailed information on the nucleation and growth behavior. For diffusion-controlled growth, one may have the following cases: 1!n!1.5 indicates the growth of particles with an appreciable initial volume; nZ1.5 indicates growth of particles with a nucleation rate close to zero; 1.5!n!2.5 reflects growth of particles with decreasing nucleation rate; nZ2.5 reflects Table 1 Avrami exponent and incubation time at different temperatures for Cu47Ti33Zr11Ni8Si1 GAP Temperature (K) 718 720 723 725 728 730 733

Incubation time (min)

Avrami exponent (0.1%Xc%0.85)

8.9 6.4 5.96 3.76 2.98 2.28 1.53

3.02 2.87 3.06 3.15 3.45 3.21 3.42

(6)

The local value of the Avrami exponent of the gasatomized powder as well as the composite powder isothermally annealed at 733 K is plotted in Fig. 9. In case of the gas-atomized powder, during the early stages of crystallization, the value is close to 2, which indicates 4.0 3.5

Avrami Exponent Nloc

723 K

720 K

725 K

728 K

730 K

733 K

ln(t-τ) (min) (b) 1.0

growth of particles with constant nucleation rate and nO2.5 pertains to the growth of small particles with an increasing nucleation rate [31]. Considering the n values we have obtained for the GAP and composite powders suggests that the crystallization is governed by an increasing nucleation rate with time. The values also suggest that with increasing annealing temperatures the rearrangement of atoms in the supercooled liquid is relatively easy. The growth of nuclei results in compositional fluctuations in the neighborhood. The composition changes result in enhanced nucleation rates at regions adjacent to the growing nuclei, which causes a chain-reaction like process leading to an increasing nucleation rate. In order to investigate the details of the crystallization process, the local Avrami exponent Nloc was calculated using the following equation [38]

3.0 2.5 2.0 1.5 1.0 GAP COMPOSITE

0.5 0.0 0

10

20

30

40

50

60

70

80

90

100

Crystallized Fractvon (%) Fig. 9. Variation of local the Avrami exponent with crystallized volume fraction for Cu47Ti33Zr11Ni8Si1 GAP and composite powder annealed at 733 K.

S. Venkataraman et al. / Intermetallics 13 (2005) 833–840

surface crystallization. With increasing crystallized volume fraction, n steadily increases and reaches a value of 3, which can be attributed to bulk crystallization. Bulk crystallization dominates when the volume fraction reaches about 5 vol.%. In this stage, the amorphous alloy crystallizes with a constant nucleation rate. In the later stages of crystallization, the value of the Avrami exponent decreases suggesting a reduced nucleation rate. This can be attributed to nucleation saturation and three-dimensional growth of crystalline nuclei leading to subsequent crystal impingement [38,39]. The observed increase in the value of the Avrami exponent in the last stage of crystallization may be due to errors derived from the isothermal DSC results. For the ball milled composite powder, the bulk crystallization stage dominates over the whole range of transformation. The local value of the Avrami exponent remains constant around 3. The increase in Nloc towards the end of the crystallization is due to increased nucleation rate in the remaining amorphous pockets. The reaction rate constant k is a function of annealing temperature and assuming it to be described by an Arrhenius equation k Z k0 ½KEc =RT;

839

for the glassy powders and the composite powders. The addition of Cu particles and ball milling up to 20 h does not seem to have any influence on the crystallization kinetics. 3.5. Interfaces between matrix and particles The presence of nanocrystalline Cu particles does not significantly change the crystallization behavior upon constant rate heating. Isothermal annealing experiments reveal that the incubation behavior of the crystallization of composite samples is not influenced to a great extent by the presence of Cu nanocrystals though the transformation to crystalline state is accelerated. Microstructure of the gas atomized powder and the composites were studied using a LEO GEMINI high-resolution scanning electron microscope in order to get a better insight of the nucleation and crystallization process. Fig. 11(a) shows the morphology of the crystallization product in the GAP, which was annealed at 723 K for 100 min. The crystallized phase is spherical and uniformly distributed throughout the volume. The spherical nature of the crystallized grains also implies a directionindependent crystal growth velocity [35]. A uniform

(7)

where k0 is a constant and Ec is the apparent activation energy for crystallization. Fig. 10 shows the plot of ln k vs. 1/T, which also yields a straight line. The activation energy for crystallization is calculated as 3.69G0.2 and 3.78G 0.1 eV for the GAP and the composite powder, respectively. These values are similar to those obtained by isochronal annealing as revealed by Kissinger analysis (3.53G0.12 and 3.59G0.06 eV) for the particle-free and the particlecontaining powder, respectively. The results indicate that both the Kissinger analysis and JMA approach can be successfully employed to predict the activation energy for the first order transition. In addition they are nearly similar –0.8 GAP COMPOSITE

–1.0 –1.2 –1.4

ln k

–1.6 –1.8 –2.0 –2.2 –2.4

EGAP

–2.6

ECOMPOSITE : 3.78 ± 0.10 eV

–2.8 1.360

1.365

: 3.69 ± 0.20 eV

1.370

1.375

1.380

1000/T

1.385

1.390

1.395

(K–1)

Fig. 10. ln k vs. 1/T for Cu47Ti33Zr11Ni8Si1 GAP as well as composite powder.

Fig. 11. High Resolution SEM Image of (a) Cu47Ti33Zr11Ni8Si1GAP annealed at 723 K for 100 min, (b) composite annealed at 723 K for 100 min.

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S. Venkataraman et al. / Intermetallics 13 (2005) 833–840

crystallite size is expected when the material crystallizes via heterogeneous nucleation [37,40]. In contrast, a non-uniform crystallite size suggests homogeneous nucleation. Thus, all the data support the view that in our glassy powder the crystalline phase nucleates homogeneously at an increasing rate and grows linearly in three dimensions. The microstructure of the composite annealed at 723 K for 100 min is shown in Fig. 11(b). It also shows a similar microstructure when compared with the as received powder. There are no reaction layers formed. Also the DSC studies do indicate that the addition of Cu has no detrimental effect on the extent of supercooled liquid region. In fact, the Cu addition increases the extent of supercooled liquid region by causing a slight decrease in the glass transition temperature. Hence there is a strong possibility that the crystallization product is formed homogenously. Although heterogenous nucleation at the interfaces cannot be completely ruled out, even if it occurs, it has no effect on the crystallization of the matrix and therefore the composite. Similar observations have been made for bulk metallic glass composites synthesized by copper mold casting [7]. The values of the Kissinger plots for the GAP powder and the composites are nearly the same. This finding has two consequences. The interfaces between the Cu particles and the amorphous matrix do not act as a heterogenous nucleation site, which would lead to a decrease the nucleation barrier. Secondly Cu does not change the growth kinetics and crystallization products since there is no change in crystallization temperature. 4. Conclusions Cu47Ti33Zr11Ni8Si1 metallic glass composite powders containing up to 25 vol.% Cu particles have been successfully synthesized by ball milling, which results in a uniform distribution of Cu particles in the glassy matrix. The thermal stability of the matrix is not affected upon the addition of the second phase particles. Activation energies for primary crystallization as calculated by Kissinger and JMA analyses are nearly equal for particle-free and composite powders. Crystallization occurs with increasing nucleation rate and the growth is three-dimensional and is diffusion-controlled. The addition of particles does not change the crystallization kinetics and the particles do not act as preferred heterogeneous nucleation sites. Ball milling is an effective tool for solid-state synthesis of metallic composite powders, which can be subsequently processed for producing bulk materials. Acknowledgements The authors thank B. Bartusch, K. Berger, and H. Schulze for technical assistance, and S. Deledda, G. Kumar and K. Biswas for helpful discussions. S. Venkataraman acknowledges funding by the German Research Foundation under grant no. Ec111/10-1,2.

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