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Nanocrystal development in Cu47Ti33Zr11Ni8Si1 metallic glass powders
Journal of Alloys and Compounds 415 (2006) 162–169
Nanocrystal development in Cu47Ti33Zr11Ni8Si1 metallic glass powders S. Venkataraman a,∗ , W. L¨oser b , J. Eckert c,a , T. Gemming d , C. Mickel a , P. Schubert-Bischoff e , N. Wanderka e , L. Schultz a , D.J. Sordelet f,g a
IFW Dresden, Institut f¨ur Metallische Werkstoffe, Postfach 27 00 16, D-01171 Dresden, Germany IFW Dresden, Institut f¨ur Festk¨orperforschung, Postfach 27 00 16, D-01171 Dresden, Germany c FG Physikalische Metallkunde, FB 11 Material und Geowissenschaften, Technische Universit¨ at Darmstadt, Petersenstraße 23, D-64287 Darmstadt, Germany d IFW Dresden, Institut f¨ ur Festk¨orperanalytik und Strukturforschung, Postfach 27 00 16, D-01171 Dresden, Germany e Hahn-Meitner-Institut Berlin, Glienickerstraße 100, D-14109 Berlin, Germany f Material and Engineering Physics Program, Ames Laboratory (USDOE), Iowa State University, Ames, IA 50014, USA g Department of Materials Science and Engineering, Iowa State University, Ames, IA 50014, USA b
Received 13 July 2005; received in revised form 4 August 2005; accepted 9 August 2005 Available online 12 September 2005
Abstract The crystallization of gas atomized Cu47 Ti33 Zr11 Ni8 Si1 metallic glass powders subjected to annealing at temperatures above and below the glass transition temperature has been studied by X-ray diffraction (XRD), differential scanning calorimetry (DSC) and transmission electron microscopy (TEM). Nucleation occurs heterogeneously at “quenched-in” nucleation sites at temperature below the glass transition. The resulting nanocrystals are Cu rich. However, upon annealing in the supercooled liquid region, crystallization of Cu51 Zr14 nanocrystals occurs homogeneously by a nucleation and growth mechanism. Crystallization in the supercooled liquid is diffusion-controlled. The activation energies for the first-order transformation have been calculated using isochronal and isothermal approaches and do not seem to be markedly affected by annealing. © 2005 Elsevier B.V. All rights reserved. Keywords: Amorphous materials; Scanning and electron microscopy; X-ray diffraction; Calorimetry
1. Introduction Metallic glass alloys show characteristic physical features, such as high strength, corrosion resistance and electromagnetic properties, which are significantly different from the corresponding crystalline alloys, due to the different atomic configurations [1]. The combination of low cost of materials, high glass-forming ability and good mechanical prop∗ Corresponding author at: FG Physikalische Metallkunde, FB 11 Material und Geowissenschaften, Technische Universit¨at Darmstadt, Petersenstraße 23, D-64287 Darmstadt, Germany. Tel.: +49 6151 166826; fax: +49 6151 165557. E-mail address: [email protected] (S. Venkataraman).
0925-8388/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2005.08.009
erties of the Cu-based amorphous alloys makes them suitable as new engineering materials [2,3]. Cu-based BMGs with high strength have been reported in Cu Ti Zr Ni [4,5], Cu Ti Zr Si B [6], Cu Ti Zr Ni Si Sn [7] and Cu (Zr/Hf) Ti [8] systems. Such alloys have very high values of fracture strength of over 2 GPa at room temperature [8]. Of particular interest is the Cu47 Ti33 Zr11 Ni8 Si1 metallic glass, which has a large supercooled liquid region prior to crystallization [9]. Bulk glassy rods of up to 7 mm in diameter are attainable [9]. The mechanical properties of Cu47 Ti33 Zr11 Ni8 Si1 have been investigated and this alloy with a “nanocomposite” microstructure, i.e. fine nanocrystals of up to 8–15 nm in dimensions uniformly distributed in an amorphous matrix exhibits fracture strength of 2040 MPa and
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also shows some plastic strain of 2.2%, slight work hardening prior to failure and a fracture strain of 4.4% [10]. There have been few reports on the crystallization of Cu47 Ti33 Zr11 Ni8 Si1 as well as Cu47 Ti34 Zr11 Ni8 glass. While Sordelet et al. [11] and Calin et al. [12] have studied in detail the phase evolution upon heating the metallic glass, Glade et al. [13] have studied the crystallization upon annealing the Cu47 Ti33 Zr11 Ni8 glass in the supercooled liquid region. The formation of a fine nanocrystalline structure has been attributed to amorphous phase separation in the as-prepared as well as isothermally annealed states [13]. Miller et al. [14] have recently shown that also the Cu47 Ti33 Zr11 Ni8 Si1 metallic glass undergoes amorphous phase separation prior to nucleation and growth of a crystalline phase. However, recent work by the present authors on the crystallization of Cu47 Ti33 Zr11 Ni8 Si1 glassy powders prepared by gas atomization has given no evidence for amorphous phase separation [15] but the nanocrystalline microstructure observed has been explained on the basis of primary crystallization. For distinguishing between crystallization reactions involving only growth, or nucleation and growth of fine crystallites, differential scanning calorimetry (DSC) can offer an effective diagnostic technique during both continuous heating and isothermal scans [16,17]. The focus of the current study is the analysis of the nanocrystal formation and the interpretation of the associated DSC patterns as well as the crystallization kinetics associated with the primary crystallization of Cu47 Ti33 Zr11 Ni8 Si1 glassy powders subjected to thermal treatment at temperatures below and above the glass transition temperature.
2. Experimental Cu47 Ti33 Zr11 Ni8 Si1 powders were produced by highpressure Ar gas atomization. The details of the powder synthesis have been reported elsewhere [15]. Annealing of the as-atomized powder was accomplished at temperatures below the glass transition temperature as well as in the supercooled liquid region, i.e. the temperature interval between the onset of the glass transition and the onset of crystallization. The isothermal annealing of the powder samples at temperatures lower than the glass transition temperature was done in a vacuum furnace operated at 1 × 10−3 bar. Prior to annealing, the powders were sealed in quartz tubes. The isothermal treatments of the as-atomized powder at temperatures above the glass transition temperature, i.e. in the supercooled liquid region, were done using a PerkinElmer DSC 7 differential scanning calorimeter (DSC) under high purity flowing Ar. Structural characterization was carried out by X-ray diffraction (XRD) using a Philips PW 1050 diffractometer (Co K␣ radiation). Calorimetric measurements of the Cu47 Ti33 Zr11 Ni8 Si1 gas atomized powder in the as-prepared as well as in the annealed condition were also carried out using the DSC. The isochronal DSC measurements were made at heating rates ranging from 10 to
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80 K/min. For the isothermal DSC measurements, the samples were first heated at a heating rate of either 20 K/min or 40 K/min to the annealing temperature and held for a certain period of time. The calibration of the DSC was done using zinc and indium standards, giving an experimental error for the temperature and the enthalpy of about 1 K and 0.5 J/g, respectively. For each individual measurement, two successive runs were performed with the second one serving as baseline. Microstructural characterization was carried out using a LEO Gemini 1531 field emission gun high-resolution scanning electron microscope (HRSEM). Transmission electron microscopy (TEM) investigations and energy dispersive X-ray analysis (EDX) were performed on powder samples using a JEOL 2000 FX TEM operated at 200 kV as well with a Tecnai F30 (FEI) high-resolution analytical TEM (HRTEM).
3. Results 3.1. As-prepared Cu47 Ti33 Zr11 Ni8 Si1 powder Fig. 1a shows the XRD patterns of the gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powder in the as-prepared state. The asatomized received powder exhibits the typical broad diffuse maxima with no traces of crystalline phases, characteristic of an amorphous phase. Fig. 2a shows the continuous heating DSC curve recorded at 20 K/min for this powder. The DSC trace exhibits an endothermic event, characteristic of the glass transition, prior to the onset of crystallization. The onset temperatures of the glass transition, Tg as well as the crystallization, Tx at 20 K/min are 691 K and 752 K, respectively. The width of the supercooled liquid region, T = Tx – Tg , is 61 K. These values are in agreement with the thermal stability data of bulk Cu47 Ti33 Zr11 Ni8 Si1 metallic glass reported in earlier studies [9,12]. A closer look at Fig. 2a reveals that there are two exothermic events for the Cu47 Ti33 Zr11 Ni8 Si1
Fig. 1. X-ray diffraction patterns (Co K␣ radiation) for the gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powders: (a) as-prepared; (b) isothermally annealed for 720 min at 623 K; (c) isothermally annealed for 60 min at 703 K; and (d) isothermally annealed for 100 min at 723 K.
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powder. The bright field image is featureless, typical of an amorphous material. The inset in Fig. 3a shows a diffuse electron diffraction halo, which is characteristic for the glassy state. 3.2. Effect of thermal treatment of Cu47 Ti33 Zr11 Ni8 Si1 powder at temperatures T < Tg
Fig. 2. DSC scan (heating rate 20 K/min) of the gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powders: (a) as-prepared; (b) isothermally annealed for 720 min at 623 K; (c) isothermally annealed for 60 min at 703 K; and (d) isothermally annealed for 100 min at 723 K.
powder upon isochronal heating up to 848 K indicating two distinct crystallization processes of the alloy. In fact, in total, four exothermic events are predicted to be associated with the nucleation and growth of crystalline phases from the amorphous precursor with two more crystallization exotherms appearing upon performing a constant heating rate DSC scan at temperatures above 848 K [11]. However, this study will focus on studying the crystallization occurring due to the appearance of the first exothermic event. The transformation enthalpy (H1 ), obtained as the area of the first exothermic peak in Fig. 2a was estimated as 27 ± 2 J/g. Fig. 3a shows the TEM image of the Cu47 Ti33 Zr11 Ni8 Si1 gas atomized
Fig. 1b shows the XRD pattern of the gas-atomized powder isothermally treated at 623 K for 720 min. The pattern does not show any crystalline reflections. It only consists of broad maxima and no appreciable diffraction peaks corresponding to crystalline phases are detected, indicating a fully amorphous nature of the powder within the resolution limit of the XRD instrument. Fig. 2b shows the corresponding DSC curve of this annealed powder. The annealing treatment has a marked influence on the isochronal curve. This is clearly visible on a magnified scale in Fig. 4a. A broad endothermic event starting at about 560 K as well as an exotherm having a peak value of 715 K are observed prior to the main crystallization exotherms, which were observed for the as-prepared powder. To find out the nature of the small exotherm, isothermal DSC studies were conducted on the samples pretreated at 623 K for 720 min. The isothermal treatment were carried out at temperatures of 663 K and 683 K. The isothermal treatment consisted of heating the powder at a heating rate of 20 K/min up to the isothermal treatment temperature holding at that temperature for 75 min. Subsequently, the sample was cooled at 100 K/min and then an isothermal run was made which serves as baseline for the first run. A monotonically
Fig. 3. Bright-field TEM images and corresponding electron diffraction patterns of: (a) as-prepared Cu47 Ti33 Zr11 Ni8 Si1 powder; (b) Cu47 Ti33 Zr11 Ni8 Si1 powder annealed at 623 K for 720 min; (c) Cu47 Ti33 Zr11 Ni8 Si1 powder heated till 785 K; and (d) HRSEM image of Cu47 Ti33 Zr11 Ni8 Si1 powder annealed at 723 K for 100 min.
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Fig. 4. DSC scan (heating rate 20 K/min) of the gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powders: (a) pre-annealed at 623 K for 720 min; (b) after second annealing at 663 K for 75 min of the pre-annealed powder; and (c) after second annealing at 683 K for 75 min of the pre-annealed powder.
decreasing signal is observed for the isothermal scans for the pretreated powder, which was isothermally annealed at 663 K and 683 K as shown in Fig. 5a and b The isochronal DSC scans after these isothermal treatments, shown in Fig. 4b and c reveal the completion of the exothermic transformation, which was observed prior to the main exotherms. The bright field TEM image of the Cu47 Ti33 Zr11 Ni8 Si1 gas atomized powder annealed at 623 K for 720 min (Fig. 3b) reveals the presence of fine nanocrystals having dimensions of up to 8 nm in an amorphous matrix. The selected area diffraction pattern (SADP) shown as inset in Fig. 3b consists of several ring patterns superimposed on a diffuse halo, which indicates the coexistence of nanocrystalline and amorphous phase. Highresolution TEM was employed for further microstructual studies. Fig. 6 shows the HRTEM image of the gas atomized powder annealed at 623 K for 720 min. The composition are found to be Cu enriched and the composition of one nanocrystal is measured as 60 ± 1 at.% Cu, 23 ± 1 at.% Ti, 10 ± 1 at.% Zr and 7 ± 1 at.% Ni. EDX analysis of the amor-
Fig. 5. Isothermal DSC scans for the gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powders: (a) annealing at 663 K of powder pre-annealed for 720 min at 623 K; (b) annealing at 683 K of powder pre-annealed for 720 min at 623 K; and (c) annealing at 703 K of as-prepared powder.
Fig. 6. Bright field HREM image of Cu47 Ti33 Zr11 Ni8 Si1 powder annealed at 623 K for 720 min.
phous matrix is close to the nominal composition (46 ± 1 at.% Cu, 31.5 ± 1 at.% Ti, 13 ± 1 at.% Zr and 10 ± 1 at.% Ni). Exact identification of the crystal structure of these fine crystals was not possible during the investigations. 3.3. Effect of thermal treatment of Cu47 Ti33 Zr11 Ni8 Si1 powder at temperature T > Tg Fig. 1c shows the XRD pattern of the gas-atomized powder annealed at 703 K for 60 min. The XRD peak around 2θ = 49◦ is somewhat sharper and narrower compared to the diffraction pattern of the as-prepared material indicating more ordering. Fig. 2c shows the DSC trace of the sample annealed at 703 K for 60 min. It is evident that there are changes in the position and area of the first exothermic event, which signifies the initiation of crystallization of the amorphous powder. However, there are no changes in the peak position for the second exothermic event. Fig. 5c shows the isothermal DSC scan of the as-atomized powder annealed at 703 K for 120 min. There is a steep decrease in the heat flow value after about 15 min of annealing. The curve displays the characteristic of a bell shape, which signifies the starting of an exothermic event. To see this effect clearly, isothermal DSC investigations of as-prepared gas atomized powders were done at temperature between 725 K and 733 K (in the supercooled liquid region) and the corresponding plots are shown in Fig. 7. Unlike the DSC traces shown in Fig. 5a and b, all the DSC traces reveal clear bell shaped signals. Additionally, all the isothermal DSC traces display an incubation time that increases with decreasing annealing time. These curves present only one exothermic peak, with a transformation enthalpy of 25 ± 2 J/g, which is close to that obtained for the first crystallization peak of isochronal heating as shown in Fig. 2a. Fig. 1d shows the
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Fig. 7. Isothermal DSC trace (after heating with 40 K/min to the annealing temperature) of as-prepared Cu47 Ti33 Zr11 Ni8 Si1 powder.
Fig. 8. Kissinger plot for evaluation of activation energy of gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powder.
XRD pattern of the glassy powder annealed at 723 K for 100 min. Crystalline reflections are clearly visible and can be indexed to the tetragonal ␥-CuTi phase as well as the hexagonal Cu51 Zr14 phase. It is also clear from the XRD pattern that a substantial portion of the powder is still amorphous. This is also corroborated by the DSC in Fig. 2d, which shows an endothermic glass transition succeeded by one exothermic event related to the crystallization of the remaining amorphous phase. 3.4. Crystallization kinetics of Cu47 Ti33 Zr11 Ni8 Si1 From the shift of the crystallization temperature in isochronal DSC scans at different heating rates (Φ = 10–80 K/min), one can deduce the activation energy for crystallization according to the Kissinger analysis [18]: Φ EP ln =− + c, (1) RTP TP2 where Φ is the heating rate, Tp is the exothermic peak temperature and R is the gas constant. The activation energy of crystallization of an amorphous phase, defined as EP , can be obtained from a plot of ln(Φ/TP2 ) versus 1/TP , yielding an approximate straight line with a slope of −EP /R. The Kissinger plots of the as-prepared and heat treated powders are shown in Fig. 8. For the as-prepared powder the EP = 3.53 ± 0.12 eV. Annealing at a temperature below Tg causes a small increase in the value to EP = 3.63 ± 0.13 eV. Upon annealing at temperatures above Tg the EP = 3.89 ± 0.12 eV is found. For isothermal DSC investigations in the supercooled liquid region, the activation energies were deduced from an Arrehnius relationship [19]: ln(τ0.5 ) =
Ec + c, TRiso
(2)
where τ 0.5 is the time for 50% transformation, Ec is the activation energy, R is the gas constant, and Tiso is the anneal-
Fig. 9. Isothermal activation energy of gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powder.
ing temperature. A plot of ln(τ 0.5 ) versus 1/Tiso gives a straight line with slope Ec /R, as shown in Fig. 9. The activation energy Ec = 4.56 ± 0.11 eV is found for the as-prepared powder while Ec = 4.41 ± 0.09 eV for the powder annealed at 623 K for 720 min. The isothermal activation energy for the powder annealed at temperatures above Tg is not shown here, since accurate calculation of the crystallization time was not possible, given the fact that the some portion of the glassy powder was already partially crystallized during the pre-annealing treatment at 703 K for 60 min (Fig. 5c).
4. Discussion From the XRD and TEM results presented, it is clear that the as-prepared atomized powders are amorphous. The DSC trace also displays a characteristic glass transition prior to the onset of crystallization. Additionally, the thermal stability data also agree well with those reported earlier [9,12]. The microstructural studies of powder annealed at temperature below the Tg unambiguously show the formation of nanocrystals. This result has not been reported before for the Cu47 Ti33 Zr11 Ni8 Si1 metallic glass.
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The isothermal DSC traces shown in Fig. 5a and b display a monotonically decreasing signal corresponding to a grain growth process [16]. The transformation enthalpy (H1 ), obtained as the area of the first exothermic peak in Fig. 2b, measured after annealing the powder for 720 min at 623 K was estimated as 26.7 ± 2 J/g. This value is quite close to the heat release measured for the first exothermic peak for the asprepared powder (27 ± 2 J/g). Thus, the possibility that the monotonically decreasing signal could be the tail of a hidden transformation can be excluded. Additionally, cooling the fully transformed sample to room temperature and rerunning the temperature program for the test gave a perfectly horizontal isothermal line, ruling out that the possibility that the signal is instrumental. The observation of a monotonically decreasing signal is similar to the growth of nanocrystals in Al Y Fe, Al Sm [20,21] and in Al Ni Ce metallic glasses [22]. Because of the multi-component nature of the glass and the negative heats of mixing of all of the metallic pairs other than Cu Ni [23], it is quite possible that there could be shortrange order domains in the liquid state between Ti Ni, Ni Zr and Cu Zr pairs. During quenching the liquid remains in a metastable equilibrium below Tg and the atoms are frozen in their liquid configuration, i.e. short-range order (quenchedin nuclei) is retained in the amorphous phase. It has been shown earlier that a significant population of clusters could be formed upon quenching to a glassy state [24]. During rapid quenching, the rate of development of clusters that may lead to crystallization by reaching a critical size may not reach a steady state for all cluster sizes [25]. This can also lead to transient effects in crystallization [24]. Annealing at T < Tg can stabilize some of such clusters to become effective nuclei. Upon annealing, these quenched-in nuclei might grow giving rise to the decaying signal observed in Fig. 5a and b. The continuously decreasing signal could also be due to changes in the short-range order (SRO) prior to crystallization caused due to segregation and clustering of one or more elements as has been reported earlier for the Fe Si B Nb Cu alloys [26,27]. A monotonically decreasing signal also implies that the crystallization proceeds with a low-energy barrier not consistent with creation of critical nuclei. This energy barrier is lower than in alloys crystallizing by nucleation and growth [28]. This suggests the possibility that a very-high density of quenched in embryos/nuclei might be present in the as prepared powder, which might subsequently grow. Thus, our studies clearly shows that nucleation at temperatures below Tg does not cause premature crystallization but that instead the nanocrystals grow from “quenched-in” nuclei. On the other hand, our XRD (Fig. 1d) and microstructural studies (Fig. 3d) on Cu47 Ti33 Zr11 Ni8 Si1 metallic glass powders annealed at 723 K for 100 min (i.e. in the supercooled liquid region) reveal the formation of a homogeneously nucleated crystalline product. This data is also corroborated by the TEM investigation of the glassy powder heated up to 785 K (completion of first exothermic peak), which is shown in Fig. 3c. The SAD pattern clearly proves
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the formation of a hexagonal Cu51 Zr14 phase. It is clear from the isothermal DSC traces (Fig. 7) that the crystallization of the Cu51 Zr14 phase occurs by nucleation and growth. As has been shown by Chen and Spaepen [16] that the isothermal calorimetric signal for a nucleation and growth process is an exothermic peak with a maximum at non-zero time. Hence, above the glass transition temperature, crystallization occurs by homogeneous nucleation. The isothermal kinetics of Cu47 Ti33 Zr11 Ni8 Si1 gas-atomized powders has been previously studied using the JMA model. It was found that the Avrami exponent, n, ranges from values of 2.8 to 3.5 in case of the as-prepared powders that were isothermally annealed in the supercooled liquid region [29]. These values suggest an increasing nucleation rate with time [30]. The Avrami constants obtained also support the possibility of crystallization by homogeneous nucleation and growth, since upon homogeneous nucleation, the nucleation rate increases until a transient steady state is reached [31]. Our Kissinger analysis revealed that the apparent activation energy for the first-order transition is 3.53 ± 0.12 eV for the as-prepared Cu47 Ti33 Zr11 Ni8 Si1 metallic glass powder. This value is close to 3.69 eV reported for a Cu47 Ti33 Zr11 Ni8 Si1 metallic glass prepared by copper mold casting [9]. The activation energies for Cu47 Ti33 Zr11 Ni8 Si1 are much higher than that of Zr55 Cu30 Al10 Ni5 glasses (2.38 eV) [32], of the well-known Zr41 Ti14 Cu12.5 Ni10 Be22.5 alloys (2 eV) [33] and the (Au85 Cu15 )77 Si9 Ge14 glassy alloy (2.49 eV) [34]. However, the value of the activation energy derived from the Kissinger equation is smaller than that of 4.3 ± 0.5 eV calculated for the primary crystallization of the Cu60 Ti20 Zr20 alloy [35]. The apparent activation energy is also not seriously affected by annealing. Hence, the presence of nanocrystals formed during annealing at temperatures below and above the glass transition do not seem to have a major effect on the crystallization of the remaining amorphous matrix. The isothermal activation energy value for the asprepared gas atomized powder is 4.56 ± 0.11 eV. This value is close to the isothermal activation energy calculated for Cu47 Ti34 Zr11 Ni8 (4.28 ± 0.11 eV) [13]. Annealing the gas atomized powder at 623 K for 720 min leads to only a marginal decrease to 4.41 ± 0.09 eV. Even though it is quite surprising that there is a large difference in activation energies calculated using the isochronal and the isothermal approaches, such deviations have been found already in case of Zr69.5 Cu12 Ni11 Al7.5 glassy alloy where the activation energy calculated using an Arrehnius type relation exceeds the value calculated by the Kissinger approach by 0.92 eV [36]. The isothermal treatments above Tg only comprise the small temperature interval (718–733 K). It is obvious that activation energies of transformation processes involving mass transport and generation of nuclei may be different from those in the wider temperature range of isochronal heating for the Kissinger analysis. Further crystallization studies are necessary to get a perfect explanation for the observed differences.
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Table 1 Enthalpies of mixing for the binary systems of the elements Cu, Ti, Zr, Ni and Si (from [23]) Hmix (kJ/mol)
Cu
Ti
Zr
Ni
Si
Cu Ti Zr Ni Si
0 – – – –
−9 0 – – –
−23 0 0 – –
+4 −35 −49 0 –
– −49 −67 −23 0
Crystallization studies on bulk Cu47 Ti34 Zr11 Ni8 metallic glass have shown that the nucleation of nanocrystals is diffusion controlled process, since the activation energy from the isothermal DSC experiments was calculated to be 4.28 ± 0.11 eV [13]. This value is close to the value attributed to titanium diffusion in Zr41.2 Ti13.8 Cu12.5 Ni10 Be22.5 (4.09 ± 0.76 eV) [37,38]. Our Ec values are close to these values and suggest that bulk crystallization in Cu47 Ti33 Zr11 Ni8 Si1 gas atomized powders in the supercooled liquid may also be dominated by Ti diffusion. Diffusion studies in amorphous alloys have shown that the late transition metal (Ni) diffuses two to six orders of magnitude faster than the early transition metal (Zr) at the same ˚ relative temperature [39]. The large size of Zr atoms (3.17 A) ˚ constituting the metallic to the other elements (2.35–2.89 A) glass limits their long-range diffusion for both nucleation and growth [40]. Previous reports suggest that the crystallization in Cu47 Ti34 Zr11 Ni8 [13], Cu47 Ti33 Zr11 Ni8 Si1 [14], as well as Zr41.2 Ti13.8 Cu12.5 Ni10 Be22.5 [41] metallic glasses is preceded by amorphous phase separation. This was believed be the reason for the formation of a fine nanocrystalline microstructure. However, a recent study on the crystallization of the Zr41.2 Ti13.8 Cu12.5 Ni10 Be22.5 glass clearly shows that there is no phase separation and the observed nanocrystalline structure can be explained by primary crystallization of an icosahedral phase [42]. In case of Cu47 Ti34 Zr11 Ni8 and Cu47 Ti33 Zr11 Ni8 Si1 metallic glasses, decomposition into copper-enriched and titanium-enriched regions was observed in the as-prepared and isothermally annealed states [13,14]. Face centred cubic nanocrystals nucleate and grow in within the amorphous matrix and the composition of one nanocrystal was found to be 80 at.% Ti and 20 at.% Zr [13]. Our results show the presence of Cu-rich nanocrystals upon annealing of the atomized powder at 623 K for 720 min. However, this does not exclude the formation of nanocrystals of other compositions. Our microstructural studies show no evidence for possible amorphous-phase separation. The enthalpy of mixing of all different pairs of elements for liquid binary systems constituting our metallic glass are shown in Table 1. Zr and Ti form an ideal solution. Ni and Cu tend to form pairs with Ti and Zr. Although the mixing enthalpy of Cu and Ni is slightly positive (+4 kJ/mol), liquid-phase separation is not likely. However, it cannot be completely ruled out. This positive enthalpy of mixing might also have an influence on phase separation in the
amorphous matrix prior to the nucleation and growth of crystalline phases. In fact, in the multicomponent metallic glasses showing clear two-phase amorphous structures, which have been reported recently, i.e. La27.5 Zr27.5 Cu10 Ni10 Al25 [43] and Ni58.5 Nb20.25 Y21.25 [44], the extremely large value of positive heat of mixing of the principal alloying elements La Zr (+13 kJ/mol) and Nb Y (+30 kJ/mol), respectively, is responsible for the amorphous phase separation. The reduced glass transition temperature (Trg = Tg /Tl ), i.e. the ratio of the glass transition temperature to the liquidus temperature, is found out to be 0.60 for the Cu47 Ti33 Zr11 Ni8 Si1 alloy. This value is very high and is typical for a good glass former [42]. It is expected that in an alloy with such a Trg , the nucleation of crystals will be easily suppressed during quenching due to a combination of driving force of nucleation and melt stability. Our XRD results indicate the appearance of crystalline reflections of ␥-CuTi as the well as the Cu51 Zr14 phase after completion of the first exothermic DSC event. It has been shown that massive crystallization of Cu47 Ti33 Zr11 Ni8 Si1 gas-atomized powders occurs by primary crystallization of Cu51 Zr14 nanocrystals while formation of ␥-CuTi occurs by growth of quenched in nuclei [15]. Kinetically, the favored crystal composition for nucleation and growth is close to that of the glass since this minimizes the atomic transport required for the transformation [34]. However, the thermodynamic driving force may not favor or even permit such a reaction. For the composition Cu47 Ti33 Zr11 Ni8 Si1 , equiatomic binary ␥-CuTi crystals are easier to form. The enthalpy of mixing for Cu Ti is −9 kJ/mol and that for Cu Zr is −23 kJ/mol [23]. Though the formation of the Cu Zr phase is more likely easier since the enthalpy of mixing is promoted by the more negative enthalpy of mixing; however, because Zr is a slow diffuser, this might be kinetically inhibited.
5. Conclusion The effect of thermal treatment at temperatures below (623 K for 720 min) and above (703 K for 60 min), the experimentally determined glass transition temperature (Tg = 691 K) on the crystallization behavior of Cu47 Ti33 Zr11 Ni8 Si1 metallic glass powders synthesized by gas atomization has been studied. Microstrucural studies clearly show the presence of nanocrystals formed in both temperature regimes. The nucleation of nanocrystals at temperatures below Tg does not seem to cause premature crystallization. DSC studies provide evidence that the formation of nanocrystals at temperatures below the glass transition occurs heterogeneously by the growth of “quenched-in nuclei” while the formation of nanocrystals at temperatures above the glass transition occurs homogeneously by a nucleation and growth mechanism. The nucleation of nanocrystals at temperatures in the supercooled liquid region is a diffusion-controlled process. Annealing of the gas-atomized powder does not have a major influence on the isochronal and isothermal activation
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energies for the crystallization of the remaining amorphous phase.
Acknowledgements The authors thank B. Bartusch and S. Scheider for technical assistance and S. Scudino and M. Stoica for stimulating discussions. Funding by the German Research Foundation under grant no. Ec 111/10-1,2 is gratefully acknowledged.
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Nanocrystal development in Cu47Ti33Zr11Ni8Si1 metallic glass powders S. Venkataraman a,∗ , W. L¨oser b , J. Eckert c,a , T. Gemming d , C. Mickel a , P. Schubert-Bischoff e , N. Wanderka e , L. Schultz a , D.J. Sordelet f,g a
IFW Dresden, Institut f¨ur Metallische Werkstoffe, Postfach 27 00 16, D-01171 Dresden, Germany IFW Dresden, Institut f¨ur Festk¨orperforschung, Postfach 27 00 16, D-01171 Dresden, Germany c FG Physikalische Metallkunde, FB 11 Material und Geowissenschaften, Technische Universit¨ at Darmstadt, Petersenstraße 23, D-64287 Darmstadt, Germany d IFW Dresden, Institut f¨ ur Festk¨orperanalytik und Strukturforschung, Postfach 27 00 16, D-01171 Dresden, Germany e Hahn-Meitner-Institut Berlin, Glienickerstraße 100, D-14109 Berlin, Germany f Material and Engineering Physics Program, Ames Laboratory (USDOE), Iowa State University, Ames, IA 50014, USA g Department of Materials Science and Engineering, Iowa State University, Ames, IA 50014, USA b
Received 13 July 2005; received in revised form 4 August 2005; accepted 9 August 2005 Available online 12 September 2005
Abstract The crystallization of gas atomized Cu47 Ti33 Zr11 Ni8 Si1 metallic glass powders subjected to annealing at temperatures above and below the glass transition temperature has been studied by X-ray diffraction (XRD), differential scanning calorimetry (DSC) and transmission electron microscopy (TEM). Nucleation occurs heterogeneously at “quenched-in” nucleation sites at temperature below the glass transition. The resulting nanocrystals are Cu rich. However, upon annealing in the supercooled liquid region, crystallization of Cu51 Zr14 nanocrystals occurs homogeneously by a nucleation and growth mechanism. Crystallization in the supercooled liquid is diffusion-controlled. The activation energies for the first-order transformation have been calculated using isochronal and isothermal approaches and do not seem to be markedly affected by annealing. © 2005 Elsevier B.V. All rights reserved. Keywords: Amorphous materials; Scanning and electron microscopy; X-ray diffraction; Calorimetry
1. Introduction Metallic glass alloys show characteristic physical features, such as high strength, corrosion resistance and electromagnetic properties, which are significantly different from the corresponding crystalline alloys, due to the different atomic configurations [1]. The combination of low cost of materials, high glass-forming ability and good mechanical prop∗ Corresponding author at: FG Physikalische Metallkunde, FB 11 Material und Geowissenschaften, Technische Universit¨at Darmstadt, Petersenstraße 23, D-64287 Darmstadt, Germany. Tel.: +49 6151 166826; fax: +49 6151 165557. E-mail address: [email protected] (S. Venkataraman).
0925-8388/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2005.08.009
erties of the Cu-based amorphous alloys makes them suitable as new engineering materials [2,3]. Cu-based BMGs with high strength have been reported in Cu Ti Zr Ni [4,5], Cu Ti Zr Si B [6], Cu Ti Zr Ni Si Sn [7] and Cu (Zr/Hf) Ti [8] systems. Such alloys have very high values of fracture strength of over 2 GPa at room temperature [8]. Of particular interest is the Cu47 Ti33 Zr11 Ni8 Si1 metallic glass, which has a large supercooled liquid region prior to crystallization [9]. Bulk glassy rods of up to 7 mm in diameter are attainable [9]. The mechanical properties of Cu47 Ti33 Zr11 Ni8 Si1 have been investigated and this alloy with a “nanocomposite” microstructure, i.e. fine nanocrystals of up to 8–15 nm in dimensions uniformly distributed in an amorphous matrix exhibits fracture strength of 2040 MPa and
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also shows some plastic strain of 2.2%, slight work hardening prior to failure and a fracture strain of 4.4% [10]. There have been few reports on the crystallization of Cu47 Ti33 Zr11 Ni8 Si1 as well as Cu47 Ti34 Zr11 Ni8 glass. While Sordelet et al. [11] and Calin et al. [12] have studied in detail the phase evolution upon heating the metallic glass, Glade et al. [13] have studied the crystallization upon annealing the Cu47 Ti33 Zr11 Ni8 glass in the supercooled liquid region. The formation of a fine nanocrystalline structure has been attributed to amorphous phase separation in the as-prepared as well as isothermally annealed states [13]. Miller et al. [14] have recently shown that also the Cu47 Ti33 Zr11 Ni8 Si1 metallic glass undergoes amorphous phase separation prior to nucleation and growth of a crystalline phase. However, recent work by the present authors on the crystallization of Cu47 Ti33 Zr11 Ni8 Si1 glassy powders prepared by gas atomization has given no evidence for amorphous phase separation [15] but the nanocrystalline microstructure observed has been explained on the basis of primary crystallization. For distinguishing between crystallization reactions involving only growth, or nucleation and growth of fine crystallites, differential scanning calorimetry (DSC) can offer an effective diagnostic technique during both continuous heating and isothermal scans [16,17]. The focus of the current study is the analysis of the nanocrystal formation and the interpretation of the associated DSC patterns as well as the crystallization kinetics associated with the primary crystallization of Cu47 Ti33 Zr11 Ni8 Si1 glassy powders subjected to thermal treatment at temperatures below and above the glass transition temperature.
2. Experimental Cu47 Ti33 Zr11 Ni8 Si1 powders were produced by highpressure Ar gas atomization. The details of the powder synthesis have been reported elsewhere [15]. Annealing of the as-atomized powder was accomplished at temperatures below the glass transition temperature as well as in the supercooled liquid region, i.e. the temperature interval between the onset of the glass transition and the onset of crystallization. The isothermal annealing of the powder samples at temperatures lower than the glass transition temperature was done in a vacuum furnace operated at 1 × 10−3 bar. Prior to annealing, the powders were sealed in quartz tubes. The isothermal treatments of the as-atomized powder at temperatures above the glass transition temperature, i.e. in the supercooled liquid region, were done using a PerkinElmer DSC 7 differential scanning calorimeter (DSC) under high purity flowing Ar. Structural characterization was carried out by X-ray diffraction (XRD) using a Philips PW 1050 diffractometer (Co K␣ radiation). Calorimetric measurements of the Cu47 Ti33 Zr11 Ni8 Si1 gas atomized powder in the as-prepared as well as in the annealed condition were also carried out using the DSC. The isochronal DSC measurements were made at heating rates ranging from 10 to
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80 K/min. For the isothermal DSC measurements, the samples were first heated at a heating rate of either 20 K/min or 40 K/min to the annealing temperature and held for a certain period of time. The calibration of the DSC was done using zinc and indium standards, giving an experimental error for the temperature and the enthalpy of about 1 K and 0.5 J/g, respectively. For each individual measurement, two successive runs were performed with the second one serving as baseline. Microstructural characterization was carried out using a LEO Gemini 1531 field emission gun high-resolution scanning electron microscope (HRSEM). Transmission electron microscopy (TEM) investigations and energy dispersive X-ray analysis (EDX) were performed on powder samples using a JEOL 2000 FX TEM operated at 200 kV as well with a Tecnai F30 (FEI) high-resolution analytical TEM (HRTEM).
3. Results 3.1. As-prepared Cu47 Ti33 Zr11 Ni8 Si1 powder Fig. 1a shows the XRD patterns of the gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powder in the as-prepared state. The asatomized received powder exhibits the typical broad diffuse maxima with no traces of crystalline phases, characteristic of an amorphous phase. Fig. 2a shows the continuous heating DSC curve recorded at 20 K/min for this powder. The DSC trace exhibits an endothermic event, characteristic of the glass transition, prior to the onset of crystallization. The onset temperatures of the glass transition, Tg as well as the crystallization, Tx at 20 K/min are 691 K and 752 K, respectively. The width of the supercooled liquid region, T = Tx – Tg , is 61 K. These values are in agreement with the thermal stability data of bulk Cu47 Ti33 Zr11 Ni8 Si1 metallic glass reported in earlier studies [9,12]. A closer look at Fig. 2a reveals that there are two exothermic events for the Cu47 Ti33 Zr11 Ni8 Si1
Fig. 1. X-ray diffraction patterns (Co K␣ radiation) for the gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powders: (a) as-prepared; (b) isothermally annealed for 720 min at 623 K; (c) isothermally annealed for 60 min at 703 K; and (d) isothermally annealed for 100 min at 723 K.
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powder. The bright field image is featureless, typical of an amorphous material. The inset in Fig. 3a shows a diffuse electron diffraction halo, which is characteristic for the glassy state. 3.2. Effect of thermal treatment of Cu47 Ti33 Zr11 Ni8 Si1 powder at temperatures T < Tg
Fig. 2. DSC scan (heating rate 20 K/min) of the gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powders: (a) as-prepared; (b) isothermally annealed for 720 min at 623 K; (c) isothermally annealed for 60 min at 703 K; and (d) isothermally annealed for 100 min at 723 K.
powder upon isochronal heating up to 848 K indicating two distinct crystallization processes of the alloy. In fact, in total, four exothermic events are predicted to be associated with the nucleation and growth of crystalline phases from the amorphous precursor with two more crystallization exotherms appearing upon performing a constant heating rate DSC scan at temperatures above 848 K [11]. However, this study will focus on studying the crystallization occurring due to the appearance of the first exothermic event. The transformation enthalpy (H1 ), obtained as the area of the first exothermic peak in Fig. 2a was estimated as 27 ± 2 J/g. Fig. 3a shows the TEM image of the Cu47 Ti33 Zr11 Ni8 Si1 gas atomized
Fig. 1b shows the XRD pattern of the gas-atomized powder isothermally treated at 623 K for 720 min. The pattern does not show any crystalline reflections. It only consists of broad maxima and no appreciable diffraction peaks corresponding to crystalline phases are detected, indicating a fully amorphous nature of the powder within the resolution limit of the XRD instrument. Fig. 2b shows the corresponding DSC curve of this annealed powder. The annealing treatment has a marked influence on the isochronal curve. This is clearly visible on a magnified scale in Fig. 4a. A broad endothermic event starting at about 560 K as well as an exotherm having a peak value of 715 K are observed prior to the main crystallization exotherms, which were observed for the as-prepared powder. To find out the nature of the small exotherm, isothermal DSC studies were conducted on the samples pretreated at 623 K for 720 min. The isothermal treatment were carried out at temperatures of 663 K and 683 K. The isothermal treatment consisted of heating the powder at a heating rate of 20 K/min up to the isothermal treatment temperature holding at that temperature for 75 min. Subsequently, the sample was cooled at 100 K/min and then an isothermal run was made which serves as baseline for the first run. A monotonically
Fig. 3. Bright-field TEM images and corresponding electron diffraction patterns of: (a) as-prepared Cu47 Ti33 Zr11 Ni8 Si1 powder; (b) Cu47 Ti33 Zr11 Ni8 Si1 powder annealed at 623 K for 720 min; (c) Cu47 Ti33 Zr11 Ni8 Si1 powder heated till 785 K; and (d) HRSEM image of Cu47 Ti33 Zr11 Ni8 Si1 powder annealed at 723 K for 100 min.
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Fig. 4. DSC scan (heating rate 20 K/min) of the gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powders: (a) pre-annealed at 623 K for 720 min; (b) after second annealing at 663 K for 75 min of the pre-annealed powder; and (c) after second annealing at 683 K for 75 min of the pre-annealed powder.
decreasing signal is observed for the isothermal scans for the pretreated powder, which was isothermally annealed at 663 K and 683 K as shown in Fig. 5a and b The isochronal DSC scans after these isothermal treatments, shown in Fig. 4b and c reveal the completion of the exothermic transformation, which was observed prior to the main exotherms. The bright field TEM image of the Cu47 Ti33 Zr11 Ni8 Si1 gas atomized powder annealed at 623 K for 720 min (Fig. 3b) reveals the presence of fine nanocrystals having dimensions of up to 8 nm in an amorphous matrix. The selected area diffraction pattern (SADP) shown as inset in Fig. 3b consists of several ring patterns superimposed on a diffuse halo, which indicates the coexistence of nanocrystalline and amorphous phase. Highresolution TEM was employed for further microstructual studies. Fig. 6 shows the HRTEM image of the gas atomized powder annealed at 623 K for 720 min. The composition are found to be Cu enriched and the composition of one nanocrystal is measured as 60 ± 1 at.% Cu, 23 ± 1 at.% Ti, 10 ± 1 at.% Zr and 7 ± 1 at.% Ni. EDX analysis of the amor-
Fig. 5. Isothermal DSC scans for the gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powders: (a) annealing at 663 K of powder pre-annealed for 720 min at 623 K; (b) annealing at 683 K of powder pre-annealed for 720 min at 623 K; and (c) annealing at 703 K of as-prepared powder.
Fig. 6. Bright field HREM image of Cu47 Ti33 Zr11 Ni8 Si1 powder annealed at 623 K for 720 min.
phous matrix is close to the nominal composition (46 ± 1 at.% Cu, 31.5 ± 1 at.% Ti, 13 ± 1 at.% Zr and 10 ± 1 at.% Ni). Exact identification of the crystal structure of these fine crystals was not possible during the investigations. 3.3. Effect of thermal treatment of Cu47 Ti33 Zr11 Ni8 Si1 powder at temperature T > Tg Fig. 1c shows the XRD pattern of the gas-atomized powder annealed at 703 K for 60 min. The XRD peak around 2θ = 49◦ is somewhat sharper and narrower compared to the diffraction pattern of the as-prepared material indicating more ordering. Fig. 2c shows the DSC trace of the sample annealed at 703 K for 60 min. It is evident that there are changes in the position and area of the first exothermic event, which signifies the initiation of crystallization of the amorphous powder. However, there are no changes in the peak position for the second exothermic event. Fig. 5c shows the isothermal DSC scan of the as-atomized powder annealed at 703 K for 120 min. There is a steep decrease in the heat flow value after about 15 min of annealing. The curve displays the characteristic of a bell shape, which signifies the starting of an exothermic event. To see this effect clearly, isothermal DSC investigations of as-prepared gas atomized powders were done at temperature between 725 K and 733 K (in the supercooled liquid region) and the corresponding plots are shown in Fig. 7. Unlike the DSC traces shown in Fig. 5a and b, all the DSC traces reveal clear bell shaped signals. Additionally, all the isothermal DSC traces display an incubation time that increases with decreasing annealing time. These curves present only one exothermic peak, with a transformation enthalpy of 25 ± 2 J/g, which is close to that obtained for the first crystallization peak of isochronal heating as shown in Fig. 2a. Fig. 1d shows the
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Fig. 7. Isothermal DSC trace (after heating with 40 K/min to the annealing temperature) of as-prepared Cu47 Ti33 Zr11 Ni8 Si1 powder.
Fig. 8. Kissinger plot for evaluation of activation energy of gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powder.
XRD pattern of the glassy powder annealed at 723 K for 100 min. Crystalline reflections are clearly visible and can be indexed to the tetragonal ␥-CuTi phase as well as the hexagonal Cu51 Zr14 phase. It is also clear from the XRD pattern that a substantial portion of the powder is still amorphous. This is also corroborated by the DSC in Fig. 2d, which shows an endothermic glass transition succeeded by one exothermic event related to the crystallization of the remaining amorphous phase. 3.4. Crystallization kinetics of Cu47 Ti33 Zr11 Ni8 Si1 From the shift of the crystallization temperature in isochronal DSC scans at different heating rates (Φ = 10–80 K/min), one can deduce the activation energy for crystallization according to the Kissinger analysis [18]: Φ EP ln =− + c, (1) RTP TP2 where Φ is the heating rate, Tp is the exothermic peak temperature and R is the gas constant. The activation energy of crystallization of an amorphous phase, defined as EP , can be obtained from a plot of ln(Φ/TP2 ) versus 1/TP , yielding an approximate straight line with a slope of −EP /R. The Kissinger plots of the as-prepared and heat treated powders are shown in Fig. 8. For the as-prepared powder the EP = 3.53 ± 0.12 eV. Annealing at a temperature below Tg causes a small increase in the value to EP = 3.63 ± 0.13 eV. Upon annealing at temperatures above Tg the EP = 3.89 ± 0.12 eV is found. For isothermal DSC investigations in the supercooled liquid region, the activation energies were deduced from an Arrehnius relationship [19]: ln(τ0.5 ) =
Ec + c, TRiso
(2)
where τ 0.5 is the time for 50% transformation, Ec is the activation energy, R is the gas constant, and Tiso is the anneal-
Fig. 9. Isothermal activation energy of gas atomized Cu47 Ti33 Zr11 Ni8 Si1 powder.
ing temperature. A plot of ln(τ 0.5 ) versus 1/Tiso gives a straight line with slope Ec /R, as shown in Fig. 9. The activation energy Ec = 4.56 ± 0.11 eV is found for the as-prepared powder while Ec = 4.41 ± 0.09 eV for the powder annealed at 623 K for 720 min. The isothermal activation energy for the powder annealed at temperatures above Tg is not shown here, since accurate calculation of the crystallization time was not possible, given the fact that the some portion of the glassy powder was already partially crystallized during the pre-annealing treatment at 703 K for 60 min (Fig. 5c).
4. Discussion From the XRD and TEM results presented, it is clear that the as-prepared atomized powders are amorphous. The DSC trace also displays a characteristic glass transition prior to the onset of crystallization. Additionally, the thermal stability data also agree well with those reported earlier [9,12]. The microstructural studies of powder annealed at temperature below the Tg unambiguously show the formation of nanocrystals. This result has not been reported before for the Cu47 Ti33 Zr11 Ni8 Si1 metallic glass.
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The isothermal DSC traces shown in Fig. 5a and b display a monotonically decreasing signal corresponding to a grain growth process [16]. The transformation enthalpy (H1 ), obtained as the area of the first exothermic peak in Fig. 2b, measured after annealing the powder for 720 min at 623 K was estimated as 26.7 ± 2 J/g. This value is quite close to the heat release measured for the first exothermic peak for the asprepared powder (27 ± 2 J/g). Thus, the possibility that the monotonically decreasing signal could be the tail of a hidden transformation can be excluded. Additionally, cooling the fully transformed sample to room temperature and rerunning the temperature program for the test gave a perfectly horizontal isothermal line, ruling out that the possibility that the signal is instrumental. The observation of a monotonically decreasing signal is similar to the growth of nanocrystals in Al Y Fe, Al Sm [20,21] and in Al Ni Ce metallic glasses [22]. Because of the multi-component nature of the glass and the negative heats of mixing of all of the metallic pairs other than Cu Ni [23], it is quite possible that there could be shortrange order domains in the liquid state between Ti Ni, Ni Zr and Cu Zr pairs. During quenching the liquid remains in a metastable equilibrium below Tg and the atoms are frozen in their liquid configuration, i.e. short-range order (quenchedin nuclei) is retained in the amorphous phase. It has been shown earlier that a significant population of clusters could be formed upon quenching to a glassy state [24]. During rapid quenching, the rate of development of clusters that may lead to crystallization by reaching a critical size may not reach a steady state for all cluster sizes [25]. This can also lead to transient effects in crystallization [24]. Annealing at T < Tg can stabilize some of such clusters to become effective nuclei. Upon annealing, these quenched-in nuclei might grow giving rise to the decaying signal observed in Fig. 5a and b. The continuously decreasing signal could also be due to changes in the short-range order (SRO) prior to crystallization caused due to segregation and clustering of one or more elements as has been reported earlier for the Fe Si B Nb Cu alloys [26,27]. A monotonically decreasing signal also implies that the crystallization proceeds with a low-energy barrier not consistent with creation of critical nuclei. This energy barrier is lower than in alloys crystallizing by nucleation and growth [28]. This suggests the possibility that a very-high density of quenched in embryos/nuclei might be present in the as prepared powder, which might subsequently grow. Thus, our studies clearly shows that nucleation at temperatures below Tg does not cause premature crystallization but that instead the nanocrystals grow from “quenched-in” nuclei. On the other hand, our XRD (Fig. 1d) and microstructural studies (Fig. 3d) on Cu47 Ti33 Zr11 Ni8 Si1 metallic glass powders annealed at 723 K for 100 min (i.e. in the supercooled liquid region) reveal the formation of a homogeneously nucleated crystalline product. This data is also corroborated by the TEM investigation of the glassy powder heated up to 785 K (completion of first exothermic peak), which is shown in Fig. 3c. The SAD pattern clearly proves
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the formation of a hexagonal Cu51 Zr14 phase. It is clear from the isothermal DSC traces (Fig. 7) that the crystallization of the Cu51 Zr14 phase occurs by nucleation and growth. As has been shown by Chen and Spaepen [16] that the isothermal calorimetric signal for a nucleation and growth process is an exothermic peak with a maximum at non-zero time. Hence, above the glass transition temperature, crystallization occurs by homogeneous nucleation. The isothermal kinetics of Cu47 Ti33 Zr11 Ni8 Si1 gas-atomized powders has been previously studied using the JMA model. It was found that the Avrami exponent, n, ranges from values of 2.8 to 3.5 in case of the as-prepared powders that were isothermally annealed in the supercooled liquid region [29]. These values suggest an increasing nucleation rate with time [30]. The Avrami constants obtained also support the possibility of crystallization by homogeneous nucleation and growth, since upon homogeneous nucleation, the nucleation rate increases until a transient steady state is reached [31]. Our Kissinger analysis revealed that the apparent activation energy for the first-order transition is 3.53 ± 0.12 eV for the as-prepared Cu47 Ti33 Zr11 Ni8 Si1 metallic glass powder. This value is close to 3.69 eV reported for a Cu47 Ti33 Zr11 Ni8 Si1 metallic glass prepared by copper mold casting [9]. The activation energies for Cu47 Ti33 Zr11 Ni8 Si1 are much higher than that of Zr55 Cu30 Al10 Ni5 glasses (2.38 eV) [32], of the well-known Zr41 Ti14 Cu12.5 Ni10 Be22.5 alloys (2 eV) [33] and the (Au85 Cu15 )77 Si9 Ge14 glassy alloy (2.49 eV) [34]. However, the value of the activation energy derived from the Kissinger equation is smaller than that of 4.3 ± 0.5 eV calculated for the primary crystallization of the Cu60 Ti20 Zr20 alloy [35]. The apparent activation energy is also not seriously affected by annealing. Hence, the presence of nanocrystals formed during annealing at temperatures below and above the glass transition do not seem to have a major effect on the crystallization of the remaining amorphous matrix. The isothermal activation energy value for the asprepared gas atomized powder is 4.56 ± 0.11 eV. This value is close to the isothermal activation energy calculated for Cu47 Ti34 Zr11 Ni8 (4.28 ± 0.11 eV) [13]. Annealing the gas atomized powder at 623 K for 720 min leads to only a marginal decrease to 4.41 ± 0.09 eV. Even though it is quite surprising that there is a large difference in activation energies calculated using the isochronal and the isothermal approaches, such deviations have been found already in case of Zr69.5 Cu12 Ni11 Al7.5 glassy alloy where the activation energy calculated using an Arrehnius type relation exceeds the value calculated by the Kissinger approach by 0.92 eV [36]. The isothermal treatments above Tg only comprise the small temperature interval (718–733 K). It is obvious that activation energies of transformation processes involving mass transport and generation of nuclei may be different from those in the wider temperature range of isochronal heating for the Kissinger analysis. Further crystallization studies are necessary to get a perfect explanation for the observed differences.
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Table 1 Enthalpies of mixing for the binary systems of the elements Cu, Ti, Zr, Ni and Si (from [23]) Hmix (kJ/mol)
Cu
Ti
Zr
Ni
Si
Cu Ti Zr Ni Si
0 – – – –
−9 0 – – –
−23 0 0 – –
+4 −35 −49 0 –
– −49 −67 −23 0
Crystallization studies on bulk Cu47 Ti34 Zr11 Ni8 metallic glass have shown that the nucleation of nanocrystals is diffusion controlled process, since the activation energy from the isothermal DSC experiments was calculated to be 4.28 ± 0.11 eV [13]. This value is close to the value attributed to titanium diffusion in Zr41.2 Ti13.8 Cu12.5 Ni10 Be22.5 (4.09 ± 0.76 eV) [37,38]. Our Ec values are close to these values and suggest that bulk crystallization in Cu47 Ti33 Zr11 Ni8 Si1 gas atomized powders in the supercooled liquid may also be dominated by Ti diffusion. Diffusion studies in amorphous alloys have shown that the late transition metal (Ni) diffuses two to six orders of magnitude faster than the early transition metal (Zr) at the same ˚ relative temperature [39]. The large size of Zr atoms (3.17 A) ˚ constituting the metallic to the other elements (2.35–2.89 A) glass limits their long-range diffusion for both nucleation and growth [40]. Previous reports suggest that the crystallization in Cu47 Ti34 Zr11 Ni8 [13], Cu47 Ti33 Zr11 Ni8 Si1 [14], as well as Zr41.2 Ti13.8 Cu12.5 Ni10 Be22.5 [41] metallic glasses is preceded by amorphous phase separation. This was believed be the reason for the formation of a fine nanocrystalline microstructure. However, a recent study on the crystallization of the Zr41.2 Ti13.8 Cu12.5 Ni10 Be22.5 glass clearly shows that there is no phase separation and the observed nanocrystalline structure can be explained by primary crystallization of an icosahedral phase [42]. In case of Cu47 Ti34 Zr11 Ni8 and Cu47 Ti33 Zr11 Ni8 Si1 metallic glasses, decomposition into copper-enriched and titanium-enriched regions was observed in the as-prepared and isothermally annealed states [13,14]. Face centred cubic nanocrystals nucleate and grow in within the amorphous matrix and the composition of one nanocrystal was found to be 80 at.% Ti and 20 at.% Zr [13]. Our results show the presence of Cu-rich nanocrystals upon annealing of the atomized powder at 623 K for 720 min. However, this does not exclude the formation of nanocrystals of other compositions. Our microstructural studies show no evidence for possible amorphous-phase separation. The enthalpy of mixing of all different pairs of elements for liquid binary systems constituting our metallic glass are shown in Table 1. Zr and Ti form an ideal solution. Ni and Cu tend to form pairs with Ti and Zr. Although the mixing enthalpy of Cu and Ni is slightly positive (+4 kJ/mol), liquid-phase separation is not likely. However, it cannot be completely ruled out. This positive enthalpy of mixing might also have an influence on phase separation in the
amorphous matrix prior to the nucleation and growth of crystalline phases. In fact, in the multicomponent metallic glasses showing clear two-phase amorphous structures, which have been reported recently, i.e. La27.5 Zr27.5 Cu10 Ni10 Al25 [43] and Ni58.5 Nb20.25 Y21.25 [44], the extremely large value of positive heat of mixing of the principal alloying elements La Zr (+13 kJ/mol) and Nb Y (+30 kJ/mol), respectively, is responsible for the amorphous phase separation. The reduced glass transition temperature (Trg = Tg /Tl ), i.e. the ratio of the glass transition temperature to the liquidus temperature, is found out to be 0.60 for the Cu47 Ti33 Zr11 Ni8 Si1 alloy. This value is very high and is typical for a good glass former [42]. It is expected that in an alloy with such a Trg , the nucleation of crystals will be easily suppressed during quenching due to a combination of driving force of nucleation and melt stability. Our XRD results indicate the appearance of crystalline reflections of ␥-CuTi as the well as the Cu51 Zr14 phase after completion of the first exothermic DSC event. It has been shown that massive crystallization of Cu47 Ti33 Zr11 Ni8 Si1 gas-atomized powders occurs by primary crystallization of Cu51 Zr14 nanocrystals while formation of ␥-CuTi occurs by growth of quenched in nuclei [15]. Kinetically, the favored crystal composition for nucleation and growth is close to that of the glass since this minimizes the atomic transport required for the transformation [34]. However, the thermodynamic driving force may not favor or even permit such a reaction. For the composition Cu47 Ti33 Zr11 Ni8 Si1 , equiatomic binary ␥-CuTi crystals are easier to form. The enthalpy of mixing for Cu Ti is −9 kJ/mol and that for Cu Zr is −23 kJ/mol [23]. Though the formation of the Cu Zr phase is more likely easier since the enthalpy of mixing is promoted by the more negative enthalpy of mixing; however, because Zr is a slow diffuser, this might be kinetically inhibited.
5. Conclusion The effect of thermal treatment at temperatures below (623 K for 720 min) and above (703 K for 60 min), the experimentally determined glass transition temperature (Tg = 691 K) on the crystallization behavior of Cu47 Ti33 Zr11 Ni8 Si1 metallic glass powders synthesized by gas atomization has been studied. Microstrucural studies clearly show the presence of nanocrystals formed in both temperature regimes. The nucleation of nanocrystals at temperatures below Tg does not seem to cause premature crystallization. DSC studies provide evidence that the formation of nanocrystals at temperatures below the glass transition occurs heterogeneously by the growth of “quenched-in nuclei” while the formation of nanocrystals at temperatures above the glass transition occurs homogeneously by a nucleation and growth mechanism. The nucleation of nanocrystals at temperatures in the supercooled liquid region is a diffusion-controlled process. Annealing of the gas-atomized powder does not have a major influence on the isochronal and isothermal activation
S. Venkataraman et al. / Journal of Alloys and Compounds 415 (2006) 162–169
energies for the crystallization of the remaining amorphous phase.
Acknowledgements The authors thank B. Bartusch and S. Scheider for technical assistance and S. Scudino and M. Stoica for stimulating discussions. Funding by the German Research Foundation under grant no. Ec 111/10-1,2 is gratefully acknowledged.
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