Effect of preannealing on glass transition and crystallization of gas atomized Cu47Ti33Zr11Ni8Si1 metallic glass powders

Effect of preannealing on glass transition and crystallization of gas atomized Cu47Ti33Zr11Ni8Si1 metallic glass powders

Intermetallics 14 (2006) 1085–1090 www.elsevier.com/locate/intermet Effect of preannealing on glass transition and crystallization of gas atomized Cu...

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Intermetallics 14 (2006) 1085–1090 www.elsevier.com/locate/intermet

Effect of preannealing on glass transition and crystallization of gas atomized Cu47Ti33Zr11Ni8Si1 metallic glass powders S. Venkataraman a,*, J. Eckert a,b, L. Schultz b, D.J. Sordelet c,d b

a IFW Dresden, Institut fu¨r Metallische Werkstoffe, Postfach 27 00 16, D 01171 Dresden, Germany FG Physikalische Metallkunde, FB 11 Material- und Geowissenschaften, Technische Universita¨t Darmstadt, 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

Available online 28 February 2006

Abstract Differential scanning calorimetry (DSC), X-ray diffraction (XRD) and transmission electron microscopy (TEM) were used as the main method to investigate the effect of relaxation on the glass transition and crystallization of Cu47Ti33Zr11Ni8Si1 metallic glass powders. The preannealing treatments were performed at temperatures close to the experimentally determined glass transition temperature. It was found that the thermal stability is profoundly affected by preannealing since the crystallization temperature is strongly influenced by preannealing and decreases with increase in preannealing temperature. In contrast, the annealing treatment does not change the glass transition temperature. During heat treatment at temperatures around the calorimetric glass transition temperature, Tg, the glassy powder undergoes microstructural alterations as revealed by TEM but not discernible by XRD. Fine nanocrystals of about 4–6 nm homogeneously dispersed in an amorphous matrix are observed by TEM after annealing at 698 K for 60 min. Kissinger analysis reveals that the preannealing decreases the activation energy for nanocrystallization thereby promoting partial crystallization of the glass. q 2006 Elsevier Ltd. All rights reserved. Keywords: B. Glasses, metallic; F. Calorimetry; F. Electron microscopy, transmission; F. Diffraction (electron, neutron and X-ray)

1. Introduction Glassy metallic alloys exhibit characteristic features such as high strength, corrosion resistance and electromagnetic properties, which are significantly different from their crystalline counterparts, due to the different atomic configurations [1]. Metallic glass formation is achieved by avoiding nucleation and growth of crystalline phases when cooling the alloy from the molten liquid [2]. Hence, studies on the crystallization processes of metallic glass-forming alloys are important for a better understanding of the glass formation and the thermal stability of the material. Low temperature annealing of metallic glasses causes changes in most physical properties [3–5]. Classical nucleation theory predicts that when metallic glasses crystallize at rather low temperatures near the glass transition temperature, Tg, * Corresponding author. Present address: FG Physikalische Metallkunde, FB 11 Material- und Geowissenschaften, Technische Universita¨t Darmstadt, Petersenstraße 23, D-64287 Darmstadt, Germany. Tel.: C 49 6151 166826; fax: C49 6151 165557. E-mail address: [email protected] (S. Venkataraman).

0966-9795/$ - see front matter q 2006 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2006.01.028

a coarse microstructure is obtained [6]. However, in many metallic glass-forming alloys, including Zr41.2Ti13.8Cu12.5 Ni10 Be22.5 [7,8], Mg62Cu25Y10Li3 [9], Zr52.5Cu17.9Ni14.6Al10Ti5 [10] and Zr57Cu15.4 Ni12.6Al10Nb5 [10], a fine microstructure of nanocrystals embedded in an amorphous matrix is observed. Different models have been proposed to explain the observed microstructure [11–14]. In case of Zr41Ti14Cu12.5Ni10Be22.5 the annealing-induced development of the local atomic structure (short-range order) is considered to have a significant influence on the subsequent crystallization of the alloy [15]. The crystallization behavior of Zr41Ti14Cu12.5Ni10Be22.5 has been shown to undergo phase separation prior to primary crystallization when annealed in the supercooled liquid region [8,15]. Depending on the thermal history, the crystallization of ZrTiCuNiBe glasses follows different pathways, which end up in different crystalline microstructures [16–18]. Cu47Ti34Zr11Ni8 is a bulk glass-forming alloy developed by Lin et al. [19]. An earlier study on the crystallization of this alloy reported on decomposition into titanium-enriched and copper-enriched regions prior to nucleation of a phase face centered cubic [20]. As shown by Choi-Yim et al. [21], Cu47Ti33Zr11Ni8Si1 bulk metallic glass has a slightly higher glass transition temperature and onset temperature of crystallization, as well as an improved stability against crystallization

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in the super cooled liquid regime as compared to the Cu47Ti34Zr11Ni8 metallic glass. Atom probe tomography (APT) characterization of the microstructure has also suggested that Cu47Ti33Zr11Ni8Si1 undergoes amorphous phase separation into titanium-rich and copper-rich regions prior to the nucleation and growth of a crystalline phase [22]. The APT studies clearly show that the amount of phase separation is more pronounced in slowly cooled injection cast samples than in rapidly quenched specimens. This suggests that there may be a tendency for phase separation in the liquid state that becomes more pronounced with decreasing cooling rate. Additionally, it has been also shown in the same study [22] that the amount of phase separation increases with heat treatment. However, a recent study on gas atomized Cu47Ti33Zr11Ni8Si1 powders has found no evidence that would suggest that nanocrystallization is preceded by amorphous phase separation [23]. Though it cannot be completely ruled out, transmission electron microscopy (TEM) observations did not give any indication for phase decomposition. Hence, the effect of a possible decomposition on the subsequent crystallization is a still a controversial question that requires further investigations. The stability of a metallic glass is linked to the nature and the nucleation kinetics of the crystalline phase(s), which form during the heat treatment of the glass [18]. The activation energies of nucleation and growth are closely related to the local atomic structure. Studies of the crystallization kinetics can provide information for better understanding the possible effects of phase separation and structural relaxation on the subsequent crystallization event [24]. In this paper, the crystallization kinetics of gas atomized Cu47Ti33Zr11Ni8Si1 metallic glass powders under different preannealing conditions at temperatures near Tg are investigated by differential scanning calorimetry (DSC), TEM and X-ray diffraction (XRD) and the effect of preannealing on the crystallization kinetics is explored. 2. Experimental Amorphous Cu47Ti33Zr11Ni8Si1 gas-atomized powder with spherical morphology and an average particle size of 45 mm was prepared by high-pressure gas atomization. The details of powder synthesis have been reported elsewhere [25]. To investigate the effect of preannealing, the as-atomized metallic glass powder (referred to as alloy A) was annealed at 688 K for 60 min (referred to as alloy B) and at 698 K for 60 min (referred to as alloy C). The isothermal annealing of the powder samples was done in a vacuum furnace operated at 1!10K3 bar. Prior to annealing the powders were sealed in quartz tubes. Structural characterization was carried out by X-ray diffraction (XRD) using a Philips PW 1050 diffractometer (Co Ka radiation). Calorimetric measurements of the Cu47Ti33Zr11Ni8Si1 powder in the as-received state as well as after annealing were carried out in a differential scanning calorimeter (Perkin Elmer DSC 7) under flowing high purity argon. The DSC measurements were made at heating rates ranging from 10 to 80 K/min. The calibration of the DSC was

done using zinc and indium standards, giving an experimental error for temperature and enthalpy of about 1 K and 0.5 J/g, respectively. For each individual measurement, two successive runs were recorded with the second one serving as the baseline. Transmission electron microscopy (TEM) investigations were performed on ion-milled powder samples using a JEOL 2000 FX TEM operated at 200 kV and a Tecnai F30 high resolution TEM (HRTEM). 3. Results and discussion Fig. 1 shows the DSC traces for the alloys A, B and C at a heating rate of 40 K/min. All DSC traces exhibit the endothermic characteristic of a glass transition followed by exothermic crystallization reactions at higher temperatures. The glass transition temperature, Tg (defined as the onset of the endothermic event), the crystallization temperature, Tx (defined as the onset of the first exothermic event) and the supercooled liquid region, DTZTxKTg and the crystallization peak temperature, Tp1 and Tp2 (defined as the peak temperatures of the first and second exothermic events, respectively) are different for the three alloys. The measured Tg, Tx, Tp1, Tp2 and DT values are listed in Table 1. The Tg, Tx, Tp1, Tp2 and DT for alloy A are 700, 762, 767, 808 and 62 K, respectively. The values of Tg, Tx and DT for alloy A are slightly lower in comparison with the thermal stability data of Cu47Ti33Zr11Ni8Si1 cast rods reported earlier by Choi-Yim et al. [21]. However, the values are higher than that of Cu47Ti33Zr11Ni8Si1 rods reported in another study [26] and for metallic ribbons of same composition [27]. After annealing at 688 and 698 K, the values of Tx, Tp1 and Tp2 are decreased compared to the as-prepared metallic glass powder (alloy A), while Tg exhibits no significant change. Moreover, prennealing lowers the DT values for alloys B and C to 53 and 49 K, respectively. These data indicate that thermal stability of the metallic glass is markedly affected by the preannealing condition. Fig. 2 (curve a) displays the XRD pattern for alloy A. No distinct crystalline phases are detected within the sensitivity

Fig. 1. DSC scans (heating rate 40 K/min) for the gas atomized Cu47Ti33Zr11Ni8Si1 powder in different states: (a) as-prepared state; (b) annealed at 688 K for 60 min; (c) annealed at 698 K for 60 min.

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Table 1 Thermodynamic and kinetic parameters of as-prepared Cu47Ti33Zr11Ni8Si1 powder (alloy A), after annealing at 688 K for 60 min (alloy B) and after annealing at 698 K for 60 min (alloy C) Samples

Tg (K)

Tx (K)

Tp1 (K)

Tp2 (K)

DT (K)

Ex (kJ/mol)

Ep1 (kJ/mol)

Ep2 (kJ/mol)

Alloy A Alloy B Alloy C

700 701 700

762 754 749

767 761 758

808 803 802

62 53 49

339G14 328G9 299G12

340G13 349G11 331G09

328G7 341G7 326G7

Tg, Tx, Tp1 and Tp2 are measured at a heating rate of 40 K/min.

limits of the XRD instrument. The XRD patterns for alloys B and C (curves b and c in Fig. 2) also exhibit the diffuse maxima, characteristic of an amorphous phase. The amorphous nature of alloy A is also ascertained by the bright field TEM image shown in Fig. 3(a), which exhibits a ‘salt and pepper’ contrast typical for amorphous materials. The inset shown in Fig. 3(a) shows a diffraction pattern with two diffraction halos, with the second one being especially faint. Fig. 3(b) shows the TEM image of alloy B. Some changes of the microstructure have occurred, which are reflected in the diffraction pattern. Compared to the diffraction pattern of alloy A (Fig. 3(a)), the first diffuse amorphous halo becomes much narrower. The second ring is more pronounced and a third very faint ring appears. These changes suggest an enhanced short-range order upon thermal treatment and point to an onset of crystallization. The microstructure of alloy C is displayed in Fig. 3(c). The bright field image as well as the diffraction pattern clearly show the presence of nanocrystals. The rings in the diffraction pattern are well defined. The inner rings are sharper as compared to the diffraction pattern of alloy B (Fig. 3(b)). In addition, a broad but faint halo can be observed, that overlaps with the inner rings. This indicates that a substantial volume fraction of the amorphous phase still exists. The sizes of the nanocrystals are around 5 nm. The exact identification of the chemical composition is not easy due to their small dimensions and the fact that they are dispersed in a chemically complex multicomponent glassy matrix. Fig. 3(d) shows a high resolution TEM (HRTEM) bright field image of alloy C. Nanocrystals of about 4–6 nm in size can be recognized by an ordered intensity pattern caused due to the regular arrangement of atoms in these crystals. However, the exact chemical composition as well as the identification of the crystal structure of these nanocrystals was not possible due to their very small size. Fig. 4 shows the continuous heating DSC traces of the asprepared Cu47Ti33Zr11Ni8Si1 metallic glass powder (alloy A) recorded at heating rates of 10, 20, 30, 40, 60, 80 K/min. As expected, the Tg, Tx, and Tp1 and Tp2 values of the alloy shift to higher temperatures with increasing heating rate. Not only the crystallization but also the glass transition displays a dependence on the heating rate during continuous heating. Similar trends exist in the DSC traces for alloys B and C (not shown). Generally, crystallization kinetics of amorphous alloys is studied in two ways: (i) non-isothermal or dynamic crystallization at a constant linear heating rate and (ii) isothermal crystallization. Non-isothermal analysis of the crystallization kinetics has become increasingly attractive since compared

with isothermal techniques, non-isothermal experiments can be performed over a shorter time period and wider temperature range. Non-isothermal data obtained from DSC studies have been generally analyzed using four well established theoretical models, i.e. the Kissinger model [28], the Ozawa model [29], the Matusita model [30] and the Gao and Wang model [31]. However, this study shall focus on the non-isothermal devitrification kinetics of the Cu47Ti33Zr11Ni8Si1 metallic glass powder using the Kissinger model. The Kissinger method is based on the assumption that if a first order transition proceeds at a rate varying with temperature, i.e. possesses activation energy, then the position of the calorimetric DSC peak, Tp, varies with the heating rate if the other experimental conditions are maintained fixed. Although, the Kissinger analysis was not originally developed for solid-state transformations it has been shown that it is applicable [32]. The activation energy depends on the temperature dependences of the nucleation and growth rates and on any transient event, which they may exhibit. It is, therefore, difficult to interpret, but is useful to compare the stability of glasses [33]. The Kissinger plots (of crystallization) for the three alloys are shown in Fig. 5(a)–(c). The activation energy for crystallization, Ex, Ep1 and Ep2 are deduced from the slope of the ln (F/T2) versus 1/T, where F denotes the heating rate and T stands for Tx and Tp1 and Tp2, are listed in Table 1. The Ex, Ep1 and Ep2 values of the as-prepared powder (alloy A) are 339G14, 340G13 and 328G7 kJ/mol, respectively. After annealing at 688 K for 60 min, the Ex, Ep1 and Ep2 values are 328G9, 349G11 and 341G7 kJ/mol, respectively, while after annealing at 698 K for 60 min there values obtained for Ex, Ep1

Fig. 2. XRD patterns for the gas atomized Cu47Ti33Zr11Ni8Si1 powder in different states: (a) as-prepared state; (b) annealed at 688 K for 60 min; (c) annealed at 698 K for 60 min.

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Fig. 3. Bright-field TEM images and corresponding electron diffraction patterns for: (a) as-prepared Cu47Ti33Zr11Ni8Si1 powder, (b) Cu47Ti33Zr11Ni8Si1 powder annealed at 688 K for 60 min, (c) Cu47Ti33Zr11Ni8Si1 powder annealed at 698 K for 60 min and (d) HRTEM image of Cu47Ti33Zr11Ni8Si1 powder annealed at 698 K for 60 min.

and Ep2 are 299G12, 331G9 and 326G7 kJ/mol, respectively. It is evident that the values of Ex, Ep1 and Ep2 are influenced by pre-annealing and the influence of pre annealing on Ex is quite strong. It was reported earlier that the onset temperature of crystallization of an amorphous alloy refers to the temperature where the crystals begin to precipitate from the amorphous matrix, while the peak temperature corresponds to the stage where the crystals begin to collide each other [34–36]. Hence, Tx is closely linked with the nucleation process and the peak temperature is more related to the growth process. It was shown that the activation energy calculated from the onset temperature of crystallization Ex represents the activation energy for nucleation, En, and the activation energy calculated from the crystallization peak temperatures Ep1 and Ep2 resemble the activation energies for growth, Eg [34]. From our results it is clear that for alloy A the Ex and Ep1 values are nearly identical. However, after preannealing the Ex values for alloys B and C are lower than the corresponding Ep1 values. Hence, the En is smaller than Eg values. It is obvious that the higher the activation energy En and/or Eg, the smaller is the nucleation rate and/or growth rate. Accordingly, preannealing decreases the activation energy for nucleation and thereby contributes to nanocrystallization. It is well known that a major prerequisite for nanocrystallization is a high nucleation rate combined with a low growth rates [37]. Our TEM investigations show that preannealing results in the formation of a fine nanocrystalline microstructure with nanocrystals of 4–6 nm present in an amorphous matrix. The kinetic data are also proving quantitatively that the nanocrystallization is easier. Previous work on Cu47Ti33Zr11Ni8Si1 suggested amorphous phase separation to occur prior to formation of nanocrystalline microstructure [22]. Amorphous phase separation was characterized by atom probe tomography (APT) and it was reported

that both the injection cast rod (cooling rate, 4!103 K/s) as well as the splat quenched ribbon (cooling rate, 2!106 K/s) had phase separated region in the as prepared state. The phase separated regions were found to be copper-enriched and titanium-enriched. However, no evidence for such a phase separation prior to nucleation and growth of a nanocrystalline phase has been found for gas-atomized powders (cooling rate between 1!104 and 5!105 K/s [38]), of same composition [23]. Since, the gas atomized powders experienced a much higher cooling rate compared to the injection cast specimens, it is quite likely that the degree of phase separation is indeed less pronounced in the faster quenched specimens, or even phase separation may be suppressed. There have been no APT studies of gas atomized Cu47Ti33Zr11Ni8Si1 metallic glass powders so far. Hence, even though phase separation in the amorphous state cannot be deduced from HRTEM investigations, it cannot be completely ruled out.

Fig. 4. DSC traces for as-prepared Cu47Ti33Zr11Ni8Si1 powder recorded at different heating rates.

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Fig. 5. Kissinger plots corresponding to: (a) crystallization onset; (b) crystallization peak temperature (peak 1) and (c) crystallization peak temperature (peak 2).

Preannealing of the metallic glass near Tg or in the supercooled liquid region can induce structural relaxation in addition to possible phase separation and nucleation and growth of nanocrystals [24]. Fig. 3 reveals that the microstructure is markedly affected by annealing and these changes are expected to influence the kinetics of crystallization. The as-prepared metallic glass must bear structural relaxation and possible decomposition in addition to nucleation and growth of nuclei in the process of continuous heating. The nucleation of alloy A is formed by a composition fluctuation from initial homogeneous alloy with lower density of nuclei. However, bulk nucleation in B and C, which have a higher density of nuclei, are formed by heterogeneous structure as well as compositional changes induced by preannealing. It is pertinent to point that the preannealing-induced microstructural changes have an impact on the nucleation and growth mechanism and hence influence the activation energy for nucleation as well as for growth. By using the non-isothermal analysis method proposed by Matusita [30] and Matusita and Sakka [39,40], an independent value of the activation energy for nucleation as well as for growth can be calculated. A more detailed study of the nucleation and growth mechanism using the above mentioned model shall be a part of a future study.

4. Conclusions This work presents an analysis of the effect of thermal treatment at temperatures around Tg on the crystallization behavior of Cu47Ti33Zr11Ni8Si1 metallic glass powders synthesized by gas atomization. Pre-annealing strongly influences the thermal stability of the glass. Annealing at temperatures near Tg promotes the formation of nanocrystals as revealed by the TEM investigations. Kissinger analysis on the kinetics of nanocrystallization reveals that preannealing

decreases the activation energy for nanocrystallization and thereby makes it easily feasible. Acknowledgements The authors thank T. Gemming, H.S. Jeevan, C. Mickel and for help with the TEM work and S. Scudino for stimulating discussions. Funding by the German Research Foundation under grant no. Ec 111/10-1,2 is gratefully acknowledged. References [1] Luborsky LE. Amorphous metallic alloys. London: Butterworths; 1983. [2] Greer AL. Science 1995;267:1947. [3] Zhu J, Clavaguera-Mora MT, Clavaguera N. Appl Phys Lett 1997;70: 1709. [4] Van den Beukel A, Sietsma J. Mater Sci Eng 1994;A179–180:86. [5] Volkert CA, Spaepen F. Acta Metall 1989;37:1355. [6] Geyer U, Schneider S, Johnson WL, Qiu Y, Tombrello TA, Macht M-P. Phys Rev Lett 1994;65:2136. [7] Peker A, Johnson WL. Appl Phys Lett 1993;63:2343. [8] Schneider S, Thiyagarajan P, Johnson WL. Appl Phys Lett 1996;68:493. [9] Liu W, Johnson WL. J Mater Res 1996;11:2388. [10] Lo¨ffler JF, Bossuyt S, Glade SC, Johnson WL, Wagner W, Thiyagarajan P. Appl Phys Lett 2000;77:525. [11] Calin M, Ko¨ster U. Mater Sci Forum 1998;269:749. [12] Schneider S, Thiyagarajan P, Geyer U, Johnson WL. Physica 1998;B 241: 918. [13] Kelton KF. Philos Mag Lett 1998;77:337. [14] Lo¨ffler JF, Johnson WL. Appl Phys Lett 2000;76:3394. [15] Wang WH, Wei Q, Friedrich S. Phys Rev 1998;B57:8211. [16] Wei Q, Wanderka N, Schubert-Bischoff P, Macht M-P, Friedrich S. J Mater Res 2000;15:1729. [17] Macht M-P, Mechler S, Mueller M, Wanderka N. Mater Sci Forum 2002; 386–388:99. [18] Mechler S, Wanderka N, Macht M-P. Mater Sci Eng 2004;A 375–377: 355. [19] Lin XH, Johnson WL. J Appl Phys 1995;78:6514.

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