Metallic glass formation in the Cu47Ti33Zr11Ni8Si1 alloy

Metallic glass formation in the Cu47Ti33Zr11Ni8Si1 alloy

Materials Science and Engineering A 444 (2007) 257–264 Metallic glass formation in the Cu47Ti33Zr11Ni8Si1 alloy S. Venkataraman a,b,∗ , B. Bartusch b...

674KB Sizes 0 Downloads 74 Views

Materials Science and Engineering A 444 (2007) 257–264

Metallic glass formation in the Cu47Ti33Zr11Ni8Si1 alloy S. Venkataraman a,b,∗ , B. Bartusch b , C. Mickel b , K.B. Kim c , J. Das a,b , S. Scudino a,b , M. Stoica a,b , D.J. Sordelet d , J. Eckert a,b a

FG Physikalische Metallkunde, FB 11 Material-und Geowissenschaften, Technische Universit¨at Darmstadt, Petersenstraße 23, D-64287 Darmstadt, Germany b Leibniz-Institut f¨ ur Festk¨orper-und Werkstoffforschung Dresden, Helmholtzstraße 20, D-01069 Dresden, Germany c Department of Advanced Materials Engineering, Sejong University, 98 Gunja-dong, Gwangjin-gu, Seoul 143-747, South Korea d Materials and Engineering Physics Program, Ames Laboratory (USDOE), Iowa State University, Ames, Iowa, IA 50014, USA Received 5 June 2006; received in revised form 18 August 2006; accepted 25 August 2006

Abstract Phase formation in the Cu47 Ti33 Zr11 Ni8 Si1 alloy synthesized by various solidification techniques has been studied using X-ray diffraction (XRD), differential scanning calorimetry (DSC) and transmission electron microscopy (TEM) paying, in particular, attention to the TEM sample preparation. While a reduced cooling rate results in a duplex microstructure, consisting of nanocrystals in an amorphous matrix, glass formation is achieved upon employing a faster quenching rate. The difference in composition and quench temperature influence the phase transition upon heating. The essential differences observed in TEM, notwithstanding the ion-milling effects and the different pre-alloy ingots used for metallic glass synthesis, can be linked to differences in the nucleation and growth mechanism which, in turn, is influenced by the chemical short range order (CSRO) obtained upon quenching. © 2006 Elsevier B.V. All rights reserved. Keywords: Metallic glasses; Crystallization; DSC; TEM; Short range order

1. Introduction The possibility of preparing bulk metallic amorphous alloys (here we define bulk as ≥2 mm at least in one dimension) has generated great interest due to the improved mechanical properties of these materials [1–3]. Of the various families of alloys developed, Cu-based bulk alloys have been reported in Cu–Ti–Zr [4], Cu–Hf–Ti [4], Cu–Zr–Al [5,6], Cu–Hf–Al [7] ternary systems and in the Cu–Ti–Zr–Ni [8], Cu–Zr–Al–Y [5], Cu–Zr–Al–Ag [9], Cu–Hf–Ti–Nb [10] and Cu–Zr–Al–Gd [11] quaternary systems. The possibility of forming bulk glasses was also explored in the Cu–Ti–Zr–Ni–M (M = Fe, B, Si, Sn, Pb, In, Ag, Cr, Mo, W) quinary systems [11–17]. Very recently the formation of bulk amorphous alloys has been reported even in Cu–Zr binary [18,19] as well as in the Cu–Ti–Zr–Ni–Si–Nb senary systems [20]. It has been clearly shown in case of Cu–Zr [18] as well as Cu–Zr–Al alloys [21] that the glass-forming ability (GFA), meaning the synthesis of a fully amorphous phase,



Corresponding author. Tel.: +49 351 4659644; fax: +49 351 4659452. E-mail address: [email protected] (S. Venkataraman).

0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.08.089

critically depends on composition and even a composition deviation of 0.5–1 at.% can degrade the GFA. X-ray diffraction (XRD) alone fails to distinguish a true amorphous structure and a fine nanocrystalline structure [22,23] and in that sense is a necessary but not sufficient tool to reveal amorphous phase formation. Clear evidence for complete glass formation can be obtained from detailed transmission electron microscopy (TEM) investigations [24]. However, the TEM preparation conditions are an important aspect to be considered and their influence on the observation of the microstructure has been stressed in various articles addressing the possibility/non-possibility of synthesizing a true amorphous structure in multicomponent metallic glass-forming alloys [24–32]. The aim of this article is to review the various issues that have arisen in literature regarding the formation of an amorphous phase in Cu-rich glass-forming compositions by studying the phase formation in Cu47 Ti33 Zr11 Ni8 Si1 specimens that were synthesized by various routes, i.e. meltspinning, copper mold casting, as well as gas-atomization, and subsequently characterized by XRD, differential scanning calorimetry (DSC) and TEM. The focus of this study is on the processing induced variations observed in synthesis of true amorphous structures.

258

S. Venkataraman et al. / Materials Science and Engineering A 444 (2007) 257–264

2. Experimental procedure The synthesis of Cu47 Ti33 Zr11 Ni8 Si1 was achieved using the well known quenching techniques like Ar-gas-atomization, melt-spinning and copper mold casting. The resultant geometry of the products was in the powder, hereinafter referred as gas-atomized powder (GAP), ribbon and rod form, respectively. The exact experimental procedures for the powder synthesis are reported in Ref. [33] while that for the ribbons and rods are given in Ref. [27]. The pre-alloy ingot in case of the GAP was prepared using a one-step procedure (i.e. arc melting all high-purity elements). However, in case of the ribbon and rod a two-step procedure, as previously described in Ref. [27] was used. The ejection temperature in case of GAP was 1623 K while in case of melt-spinning and copper mold casting it was 1473 K. The ejection temperature was measured using an external two-color optical pyrometer. The average particle size of the powder used in this study was 20 ␮m. The ribbon had a cross section of 0.05 mm × 3 mm while the diameter of the copper mold cast rod was 4 mm. XRD measurements were done with a Philips PW 1050 diffractometer with Co K␣ radiation. Composition analysis was carried out on the as-prepared state of these materials using inductively coupled plasma optical emission spectroscopy (ICP-OES). The amount of oxygen (in weight ppm) was evaluated by carrier gas hot extraction using a Leco T6-436DR analyzer. The thermal stability was investigated by DSC at 20 K/min heating rate with a Perkin-Elmer DSC7 under a continuous flow of purified argon. 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. Additionally, some isothermal experiments were also conducted on the samples. They consisted of heating the specimens at a heating rate of 40 K/min up to 723 K and holding for 45 min. Both, in case of isothermal and isochronal treatments, two identical runs were made sequentially for each sample with the second one serving as the baseline. At least 5 DSC measurements were done for all the samples in order to confirm their reproducibility. The liquidus temperature was measured upon heating using a Netzsch DSC 404 operated at a heating rate of 20 K/min. The liquidus temperature was determined as the inflection point of the higher temperature side of the last endotherm of the heating curve. Microstructural characterization was carried out using a Philips CM 20 TEM operated at 200 kV accelerating voltage. The TEM specimens were prepared with a BAL-TEC RES 010 rapid etching system. The ion-milling was done under liquid nitrogen with 3 kV and 3 mA as the ion-milling parameters. The milling angle was 7◦ . In order to study the possible influence of TEM preparation conditions, additional samples were also prepared without liquid nitrogen. In this case, the TEM specimens were prepared by ion-milling using a Gatan precision ion mill (PIPS model 641) with an accelerating voltage of 3 kV and ion beam current of 5 mA. In case of the ribbons, the TEM samples were taken from the wheel side while in case of the rod, they were taken from the centre of the sample. Since it is well known that ion milled samples of Cu–Zr-based glass-forming alloys oxidize readily on exposure to air atmosphere [29], TEM observations were done directly

after preparation of samples to avoid observation of oxidationinduced artifacts. 3. Results 3.1. Calorimetric observations Fig. 1 (upper part) shows the constant heating rate DSC scans (20 K/min) for all the samples. The DSC traces exhibit a distinct glass transition (clearly visible in the lower part of Fig. 1), followed by a supercooled liquid region before two crystallization events occur at higher temperature. The onset temperature of the glass transition, Tg , is 691, 689 and 691 K for the GAP, the rod and the ribbon, respectively. The crystallization temperature Tx1 , defined as the onset of the first exothermic DSC event, is 752, 751 and 753 K, respectively. The supercooled liquid region, Tx = Tx1 − Tg , that corresponds to the temperature range of the stability of the supercooled liquid against crystallization, is found to be 61 K for the GAP and 62 K for both the rod and the ribbon specimens. The exothermic heat release upon the completion of the first crystallization event (Hx1 ) is 30 ± 1 J/g for the GAP while it is 33 ± 1 for the ribbon and the rod. Although the Tg , Tx1 and Hx1 are very similar for all the samples, there are essential differences in the DSC traces. The most obvious difference between the DSC traces is seen in case of the second exothermic event (Tx2 ), for the GAP vis-a-vis that of the ribbon and the rod. The second exothermic event in case of the GAP has an onset value of 787 K which is lower that that observed in case of the ribbon and rod (805 K). However, the exothermic heat release upon the completion of the second crystallization event (Hx2 ) is 27 ± 1 J/g for all the different samples. Fig. 1 (lower part) shows the enlarged low-temperature part of the DSC traces (20 K/min). Upon heating the samples, in case of the ribbon as well as the rod, the heat flow essentially remains constant up to 680 K before an increase in the endothermic heat flow signal sets in. Contrarily, in case of the GAP there is a low temperature exotherm that starts at around 490 K and has a minimum at around 660 K. The thermal stability data for all the specimens are summarized in Table 1.

Fig. 1. Constant heating rate (20 K/min) DSC scans of the different Cu47 Ti33 Zr11 Ni8 Si1 specimens.

S. Venkataraman et al. / Materials Science and Engineering A 444 (2007) 257–264

259

Table 1 Thermal stability data for the different Cu47 Ti33 Zr11 Ni8 Si1 samples by DSC at a constant heating rate of 20 K/min Specimen geometry

Tg (K)

Tx1 (K)

Tx2 (K)

T (Tx1 − Tg ) (K)

Hx1 (J/g)

Hx2 (J/g)

GAP Rod Ribbon

691 689 691

752 751 753

787 807 808

61 62 62

30 (±1) 33 (±1) 33 (±1)

27 (±1) 27 (±1) 27 (±1)

Fig. 2. Isothermal DSC traces of the different Cu47 Ti33 Zr11 Ni8 Si1 specimens. The inset shows JMA plots for 0.10 ≤ Xc ≤ 0.85.

Fig. 2 shows the isothermal DSC traces for the different samples at 723 K. The onset of crystallization as evidenced by an exothermal “bell” shaped curve begins at an earlier time for the ribbon and the rod as compared with the GAP. The time evolution of the fraction of a phase is conventionally represented by describing the kinetics of isothermal phase transformation, known as the Johnson–Mehl–Avrami (JMA) model [34]. The essence of the model can be mathematically written in the form of the JMA equation: x(t) = 1 − exp{−[k(t − τ)]n },

(1)

here, x(t) is the volume fraction of the transformed phase, t the annealing time, n the Avrami constant related to the dimensionality of nucleation and growth, k a kinetic constant and τ is the incubation time. The Avrami exponent is found to be 3.06 for the GAP, 3.42 for the ribbon and 3.54 for the rod. Isochronal DSC experiments (not shown here) were also conducted at different heating rates (10–80 K/min). The activation energy for the crystallization was calculated using the Kissinger approach [35] (as given by Eq. (2)) from the shift of the crystallization peak temperature (TP1 ) with different heating rates (ϕ):   E ϕ = ln + c, (2) 2 RTP1 TP1 here, E is the activation energy for crystallization and R is the gas constant. Fig. 3 shows the Kissinger plots for the different Cu47 Ti33 Zr11 Ni8 Si1 samples. The activation energy is found from the slope of the curve and in case of all the geometries examined, Eq. (2) is valid with the correlation coefficient ≥0.998. The activation energy for the first crystallization event is 3.53, 2.86 and 2.96 eV for the GAP, the rod and the ribbon, respectively. Fig. 4 shows the melting endotherms measured using the high-temperature DSC for all the specimens. The melting signatures for the ribbon as well as the rod are quite similar while that for the GAP is different. These alloys are at

Fig. 3. Kissinger plots for the different Cu47 Ti33 Zr11 Ni8 Si1 specimens.

260

S. Venkataraman et al. / Materials Science and Engineering A 444 (2007) 257–264

Fig. 4. High temperature DSC scans (20 K/min) of the different Cu47 Ti33 Zr11 Ni8 Si1 specimens.

off-eutectic compositions since there are at least two events in the melting traces for each individual sample, signifying a wide temperature range for melting in contrast to the melting trace for a perfect eutectic composition, which should display, a single endothermic event upon heating [36]. 3.2. Microstructural characterization 3.2.1. X-ray diffraction structure Fig. 5 displays the XRD patterns of the samples in the asprepared condition (Fig. 5(a)) and also after heating up to the completion of the first crystallization event (785 K) (Fig. 5(b)) and after completion of the second crystallization event (838 K) (Fig. 5(c)). All the patterns in Fig. 5(a) show the typical broad maxima of an amorphous structure without any indication of crystallization. Upon heating to the end of the first crystallization event (785 K) (Fig. 5(b)) there is a sharp crystalline reflection appearing at 2θ ≈ 48 which can be indexed as the ␥CuTi compound with tetragonal P4/nmm structure [37] as well as with the Cu51 Zr14 compound with hexagonal P6/m structure [38]. However, no clear distinction between these two phases can be made from the XRD patterns. Upon heating the samples to 838 K (Fig. 5(c)) numerous crystalline peaks appear in the XRD pattern and they can be unambiguously indexed as reflections from the ␥-CuTi and Cu51 Zr14 phases. However, in case of the GAP there is a leftward shift of the peak at about 2θ ≈ 47.5◦ , identified as the (2 1 3) reflection of the Cu51 Zr14 phase. 3.2.2. Transmission electron microscopy (TEM) structure Figs. 6 and 7 show the TEM bright field images as well as the corresponding selected area diffraction (SAD) patterns (shown as insets) for the different samples. In case of the rod, the TEM specimens were taken from the centre while for the ribbon the wheel side of the ribbon was studied. In case of the GAP, powders of less than 20 ␮m in size were investigated. Fig. 6 displays the TEM results for the different samples using conventional ion-milling without any cooling, while Fig. 7 shows the TEM observations for the specimens prepared using liquid nitrogen cooling. It is clearly seen that the TEM preparation conditions have a noticeable influence on the

Fig. 5. XRD patterns for the different Cu47 Ti33 Zr11 Ni8 Si1 specimens: (a) as prepared state, (b) after heating to 785 K, and (c) after heating to 838 K.

observation of the intrinsic microstructure in case of the asspun ribbon. The TEM sample for the ribbon prepared under liquid nitrogen cooling (Fig. 7) is amorphous. Contrarily, the TEM sample prepared without liquid nitrogen (Fig. 6) shows a nanocrystalline microstructure. On the other hand, the 4 mm diameter cast rod (Figs. 6 and 7) displays a nanocrystalline microstructure for both sets of TEM preparation conditions, i.e. liquid nitrogen ion-milling as well as room temperature ionmilling. The nanocrystals are indicated by arrows. The SAD pattern for the 4 mm rod in Fig. 6 shows sharpening of the second diffuse ring while the SAD pattern for the 4 mm rod in Fig. 7 displays distinct spots. In case of the GAP, a featureless microstructure, typical of an amorphous phase, is observed upon liquid nitrogen cooling as well as without liquid nitrogen cooling.

S. Venkataraman et al. / Materials Science and Engineering A 444 (2007) 257–264

261

Fig. 6. Bright field TEM images and corresponding SAD patterns (shown as insets) for the different Cu47 Ti33 Zr11 Ni8 Si1 specimens (ion-milling conditions, 3 kV, 5 mA, no liquid nitrogen cooling).

4. Discussion Table 2 shows the chemical compositions of the GAP, the rod and the ribbon as measured by ICP-OES. It can be seen that the ribbon and the rod have an actual composition that is nickel depleted when compared with the nominal composition. Contrarily, the GAP has an actual composition, which is very close to the expected one. Also given in Table 2 are the oxygen contents measured for the different samples. Apparently, the oxygen content in the GAP is higher compared to the amount of oxygen found in the rod and the ribbon. The DSC traces (Fig. 1) essentially are a manifestation of the effect of changed composition. In case of the GAP, there is a broad exothermic event visible prior to the endothermic event characteristic of glass transition as well the main exothermic event representing crystallization. A previous study by the authors has found that the origin of this exotherm is due to structural relaxation and nanocrystallization prior to bulk crystallization [33]. Additionally, the peak for the second crystallization event is shifted to a lower temperature

(787 K) instead of 805 K. X-ray diffraction results indicate that that phase formation upon the completion of the first crystallization event could be either the tetragonal ␥-CuTi phase or hexagonal Cu51 Zr14 . The Kissinger analysis reveals that indeed the activation energy values for the first crystallization event are quite different. They are 3.53 eV for the GAP but around 2.96 eV and 2.86 eV for the ribbon and the rod, respectively. The magnitude of the calculated Kissinger activation energies are a measure of the relative stability of the starting phase and imply the necessity for a higher thermal energy for the atomic movement [39,40]. Hence, it can be said that the amorphous phase in the GAP is stronger resistant to crystallization. The fact that the activation energy for the GAP is larger in spite of higher amount of oxygen suggests that either oxygen stabilizes the glassy phase or possibly does not influence glass formation. Similar results were obtained in case of NiZr2 metallic glass ribbons which were intentionally contaminated with oxygen [39]. Another possible reason for the difference in the Kissinger activation energy could be that the values correspond to different transitions. The

Table 2 Composition analysis and oxygen content for the different Cu47 Ti33 Zr11 Ni8 Si1 samples Specimen geometry

Cu (at.%)

Ti (at.%)

Zr (at.%)

Ni (at.%)

Si (at.%)

O2 (ppm)

GAP Rod Ribbon Nominal composition

46.47 48.15 48.61 47.00

32.66 34.00 33.78 33.00

11.11 11.60 11.48 11.00

9.17 5.17 4.82 8.00

0.58 1.05 1.29 1.00

2400 900 600

262

S. Venkataraman et al. / Materials Science and Engineering A 444 (2007) 257–264

Fig. 7. Bright field TEM images and corresponding SAD patterns (shown as inset) for the different Cu47 Ti33 Zr11 Ni8 Si1 specimens prepared using ion-milling (ion-milling conditions, 3 kV, 3 mA, liquid nitrogen cooling).

presence of oxygen can suppress/support the formation of stable/metastable phases upon crystallization [39,41–45], which can have different activation energies for transformation. In an earlier study on the GAP we found that the product of the first crystallization exotherm is the Cu51 Zr14 phase [33]. Another study based on synchrotron data on the cast rods has shown that the product of the first crystallization exotherm is the ␥-CuTi phase [14]. However, it needs to be mentioned that the differences in the phase formation could be attributed to the different heating procedures being adopted. In case of DSC the heating rate is 40 K/min while in case of the synchrotron investigations it was 5 K/s. It has been clearly demonstrated that the thermal events observed upon heating the Cu47 Ti33 Zr11 Ni8 Si1 metallic glass at 40 K/min in the DSC do not necessarily correspond to the formation of equilibrium phases [46]. Further, it is clearly evident from Fig. 2 that the different Cu47 Ti33 Zr11 Ni8 Si1 samples exhibit different incubation times upon annealing at 723 K. The GAP shows a better stability against crystallization (despite its higher oxygen content). The incubation time is the time interval that must elapse prior to the formation of a significant number of nuclei and depends on the embryo distribution characteristic which is retained upon quenching [47]. The Avrami values calculated from the isothermal data provides further insight into the crystallization mech-

anism [48]. Large Avrami exponents (n > 3) correspond to crystallization accompanied with an increasing nucleation rate [49]. This means crystallization proceeds by growth of the appearing new nuclei. It has been suggested that the Avrami exponent (n) is given by the following equation [47]: n = n n + ng ,

(3)

where nn is related to nucleation and ng is related to growth. Crystallization of metallic glasses in the supercooled liquid region is a diffusion controlled process [50]. Hence, the assumption that ng = 1.5 is valid [47]. Thus, a value of nn of 1.56, 2.04 and 1.92 is obtained for the GAP, the rod and the ribbon, respectively. An estimate of the nucleation rate (I) can be made using the nn values using the relation [47] I = ct nn −1 ,

(4)

where c is a constant and t is the time. It is evident that in case of GAP I ≈ t1/2 while in case of the rod and ribbon I ≈ t. It is pertinent to point out that the GAP was quenched from about 1600 K while the ribbon and the rod were quenched from about 1500 K. Hence, nucleation in the GAP, which is quenched from a higher temperature, is more difficult, requiring more time and energy [51]. Additionally, the liquidus temperature is determined to be

S. Venkataraman et al. / Materials Science and Engineering A 444 (2007) 257–264

about 1160 K for all the different specimens with slightly different compositions. Since the quench temperature for the rod and the ribbon (1473 K) is closer to the liquidus temperature (1160 K) compared with that of the GAP (1673 K) it is plausible to assume that the degree of chemical short range order (CSRO) is different in case of the GAP than in the rod and the ribbon. In fact, neutron diffraction studies in liquid PbLi alloys show a decrease in CSRO with increasing temperature above the liquidus [52]. A recent report by Schroers and co-workers [53] on strain rate dependent cooling rate effects for the Au-based metallic glass also shows that at higher processing temperature, for a given strain rate the crystalline volume fraction decreases. This also indirectly suggests that the CSRO decreases with increasing temperature. A previous study on Pd40 Ni40 P20 metallic glass quenched at different rates showed that the crystallization behavior upon linear heating in a DSC cell was different [54]. It was reasoned that inhomogeneous phosphorous distribution in the melt gave rise to differences in short range order and hence differences in the crystallization behavior [54]. Given the slight differences in the starting composition (Table 2) as well as the processing temperature it is logical to assume that the CSRO could be different. The TEM data show that the sample preparation technique has indeed an influence on the observation of the microstructure of the ribbon. In case of the rod, a nanocrystalline microstructure can be observed irrespective of the TEM preparation technique. The sharpening of the second ring, seen in the SAD pattern of the rod (Fig. 6) reveals ordering. Additionally the SAD pattern of the rod (Fig. 7) exhibits the spots which are a characteristic of a crystal. It is, however, not surprising that the SAD pattern rings exhibit a mostly diffusive nature given the predominant contribution from the amorphous phase. The formation of a composite [nanocrystals in a majority amorphous phase] can be explained on the basis of a decreased cooling rate (estimated to be ∼102 or 103 K/s in case of the rod see, Ref. [55]). However, upon TEM observations we do observe that the GAP is fully amorphous. The cooling rates experienced by the ribbon as well as the GAP are essentially of the same order (105 –106 K/s [56]). The possible reasons for the differences in the TEM observations are more likely due to a combination of factors including the composition change, differences in short range order and nucleation rate. It is pertinent to note that of all the Cu-based glass-forming alloys, the Cu60 Zr30 Ti10 composition has been most extensively studied [22,23,25,26]. The effect of TEM preparation conditions is well established for this ternary composition clearly showing that the preparation conditions in fact influence the observation of the true microstructure, especially in case of as-spun ribbons [29]. However, a closer look at the thermal stability data from these reports [22,25,26] reveals that there are essential changes in the Tg and Tx values. An additional support for this is also obtained form the results of numerous crystallization studies on the Cu60 Zr30 Ti10 glass where a bcc phase, a fcc phase, or a hexagonal phase have been reported as the primary precipitating phase [23,57–59]. Interestingly, the thermal stability data for all the above mentioned Cu60 Zr30 Ti10 compositions are different [23,57–59]. Hence, it is very likely that a slight change in composition also has a decisive influence on the development of a

263

particular short range order and hence the preferential formation of particular phase(s). Ion-milling can induce additional defects into the amorphous structure [60]. However, localized heating can even occur during low-temperature ion-milling [60]. It has been found by Sun et al. [24] that a temperature rise during milling (estimated to be 373 K at 4 keV with a milling angle less than 15◦ ) together with the radiation effects caused the crystallization of the Cu64.5 Zr35.5 glass even though it has a relatively high glass transition (757 K) and crystallization temperature (765 K). Additionally, it has been shown from electron scattering data that in case of Zr57 Ti5 Cu20 Ni8 Al10 bulk glass ion-milling results in the appearance of sharper diffraction features [61]. This suggests an increased short range ordering [61]. However, the highresolution image shows no clear crystalline features. Radiation effects like enhanced diffusion, coupled with compositional as well as structural changes can also have an influential role in the development of ordered clusters, which can form precursors for crystal formation [24,60,61]. In case of ion-milling at room temperature there is a higher possibility of defect structures created during the process also getting annihilated subsequently. However, this may not be necessarily true in case of ion-milling at liquid nitrogen temperatures since the defects once formed may possibly get frozen and given the lower temperature may not anneal out. More detailed studies are necessary to clarify this issue, especially, in metallic glasses. Summarizing, it seems that the synthesis of a fully glassy microstructure in the Cu47 Ti33 Zr11 Ni8 Si1 glass-forming alloy is crucially dependent on numerous processing factors. These include the composition, the oxygen content and the quench temperature in addition to the TEM sample preparation effects. 5. Conclusions The effect of different preparation conditions on the formation of the Cu47 Ti33 Zr11 Ni8 Si1 metallic glass has been studied. The preparation conditions and processing have a significant impact on the final composition of the solidified glassy phase. In spite of the fact that the GAP and the ribbon experienced similar cooling rates, the oxygen stabilized amorphous phase in GAP is more resistant to the effect of TEM preparation conditions. The observation of an amorphous phase seems to be more dependent on the different nucleation and growth mechanisms. Acknowledgements The authors thank J. Acker, K. Albe, R. Benz, K. Biswas, C. Brockmann, C. Fasel, M. Frey, A. Ramar, W. Xu, P. Yu and L.C. Zhang for technical assistance and stimulating discussions. Funding for this work was provided by the German Research Foundation under grant no. Ec 111/10 as well as by the European Union within the framework of the Research Training Network on “ductile metallic glass composites” (MRTN-CT2003-504692). Synthesis efforts by DJS were supported by the US Department of Energy, Basic Energy Sciences, through Iowa State University under Contract No. W-7405-ENG-82.

264

S. Venkataraman et al. / Materials Science and Engineering A 444 (2007) 257–264

References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32]

A. Inoue, Acta Mater. 48 (2000) 279–306. M. Telford, Mater. Today 7 (2004) 36–43. W.H. Wang, C. Dong, C.H. Shek, Mater. Sci. Eng. R 44 (2004) 45–89. A. Inoue, W. Zhang, T. Zhang, K. Kurosaka, Acta Mater. 49 (2001) 2645–2652. D. Xu, J. Duan, W.L. Johnson, Phys. Rev. Lett. 92 (2004) 2455041–2455044. J. Das, M.B. Tang, K.B. Kim, R. Theissmann, F. Baier, W.H. Wang, J. Eckert, Phys. Rev. Lett. 94 (2005) 205501–205504. P. Jia, H. Guo, Y. Li, J. Xu, E. Ma, Scripta Mater. 54 (2006) 2165–2168. X.H. Lin, W.L. Johnson, J. Appl. Phys. 78 (1995) 6514–6519. J.C. Oh, T. Ohkubo, Y.C. Kim, E. Fleury, K. Hono, Scripta Mater. 53 (2005) 165–169. C. Qin, W. Zhang, K. Asami, N. Ohtsu, A. Inoue, Acta Mater. 53 (2005) 3903–3911. H.M. Fu, H. Wang, H.F. Zhang, Z.Q. Hu, Scripta Mater. 55 (2006) 147–150. H. Choi-Yim, R. Busch, W.L. Johnson, J. Appl. Phys. 12 (1998) 7993–7997. E.S. Park, H.K. Lim, W.T. Kim, D.H. Kim, J. Non-Cryst. Solids 298 (2002) 15–22. M. Calin, M. Stoica, J. Eckert, A.R. Yavari, L. Schultz, Mater. Sci. Eng. A 392 (2005) 169–178. E.S. Park, W.T. Kim, D.H. Kim, Mater. Trans. 45 (2004) 2693–2696. E.S. Park, H.J. Chang, D.H. Kim, T. Ohkubo, K. Hono, Scripta Mater. 54 (2006) 1569–1573. B. Liu, L. Liu, Mater. Sci. Eng. A 415 (2006) 286–290. D. Wang, Y. Li, B.B. Sun, M.L. Sui, K. Lu, E. Ma, Appl. Phys. Lett. 84 (2004) 4029–4031. D. Xu, B. Lohwongwatana, G. Duan, W.L. Johnson, C. Garland, Acta Mater. 52 (2004) 2621–2624. E.S. Park, D.H. Kim, T. Ohkubo, K. Hono, J. Non-Cryst. Solids 351 (2005) 1232–1238. D. Wang, H. Tan, Y. Li, Acta Mater. 53 (2005) 2969–2979. J.Z. Jiang, J. Saida, H. Kato, T. Ohsuna, A. Inoue, Appl. Phys. Lett. 82 (2003) 4041–4043. J.Z. Jiang, H. Kato, T. Ohsuna, J. Saida, A. Inoue, K. Saksl, H. Franz, K. Stahl, Appl. Phys. Lett. 83 (2003) 3299–3301. B.B. Sun, Y.B. Wang, J. Wen, H. Yang, M.L. Sui, J.Q. Wang, E. Ma, Scripta Mater. 53 (2005) 805–809. E.S. Park, H.J. Chang, D.H. Kim, W.T. Kim, Y.C. Kim, N.J. Kim, Y.W. Kim, Scripta Mater. 51 (2004) 221–224. D. Nagahama, T. Ohkubo, T. Mukai, K. Hono, Mater. Trans. 46 (2005) 1264–1270. S. Venkataraman, M. Stoica, S. Scudino, C. Mickel, T. Gemming, U. Kunz, K.B. Kim, L. Schultz, J. Eckert, Scripta Mater. 54 (2006) 835–840. D. Nagahama, T. Ohkubo, K. Hono, Scripta Mater. 49 (2003) 729–734. H.J. Chang, E.S. Park, Y.C. Kim, D.H. Kim, Mater. Sci. Eng. A 406 (2005) 119–124. N. Tian, M. Ohnuma, T. Ohkubo, K. Hono, Mater. Trans. 46 (2005) 2880–2885. D.C. Hofmann, G. Duan, W.L. Johnson, Scripta Mater. 54 (2006) 1117–1122. Y.B. Wang, H.W. Yang, B.B. Sun, B. Wu, J.Q. Wang, M.L. Sui, E. Ma, Scripta Mater. 55 (2006) 469–472.

[33] S. Venkataraman, S. Scudino, J. Eckert, T. Gemming, C. Mickel, L. Schultz, D.J. Sordelet, J. Mater. Res. 21 (2006) 597–607. [34] M. Avrami, J. Chem. Phys. 7 (1939) 1103–1112. [35] H.E. Kissinger, Anal. Chem. 29 (1957) 1702–1706. [36] A.D. Pelton, Phase diagrams, in: R.W. Cahn, P. Haasen (Eds.), Physical Metallurgy, Part I, North Holland, Amsterdam, 1983, p. 336. [37] PDF #07-0114, PCPDF Version 2.2, JCPDS-International Centre for Diffraction Data (ICDD), Newton Square, PA, 2001. [38] PDF #42-1185, PCPDF Version 2.2, JCPDS-International Centre for Diffraction Data (ICDD), Newton Square, PA, 2001. [39] S. Boutet, G. Steele, M. Dikeakos, Z. Altounian, J. Appl. Phys. 84 (2001) 2441–2446. [40] Z. Altounian, J.O. Strom-Olsen, in: R.D. Shull, A. Joshi (Eds.), Thermal Analysis in Metallurgy, The Minerals, Metals and Materials Society, Warrendale, PA, 1992, p. 155. [41] L.Q. Xing, A. Mukhopadhyay, A. Buhro, W.E. Buhro, K.F. Kelton, Philos. Mag. Lett. 84 (2004) 293–302. [42] J. Eckert, N. Mattern, M. Zinkevitch, M. Siedel, Mater. Trans. JIM 39 (1998) 623–632. [43] A. Gebert, J. Eckert, L. Schultz, Acta Mater. 46 (1998) 5475–5482. [44] B.S. Murty, D.H. Ping, K. Hono, A. Inoue, Mater. Sci. Eng. A 304–306 (2001) 706–709. [45] D.J. Sordelet, X.Y. Yang, E. Rozhkova, M.F. Besser, M.J. Kramer, Appl. Phys. Lett. 83 (2003) 69–71. [46] D.J. Sordelet, M.J. Kramer, M.F. Besser, E. Rozhkova, J. Non-Cryst. Solids 290 (2001) 163–172. [47] J. Burke, Kinetics of Phase Transformation in Metals, Pergamon, Oxford, UK, 1965. [48] J.W. Christian, The Theory of Phase Transformation in Metals and Alloys, Pergamon, Oxford, UK, 1965. [49] R.D. Doherty, in: R.W. Cahn, P. Haasen (Eds.), Physical Metallurgy, vol. II, North Holland, Amsterdam, 1983. [50] U. K¨oster, U. Herold, in: H.J. Guntherodt, H. Beck (Eds.), Topic in Applied Physics, 46, Spriger Verlag, Berlin, 1981, pp. 225–257. [51] Z. Altounian, J.O. Strom-Olsen, J.L. Walter, J. Appl. Phys. 55 (1984) 1566–1571. [52] H. Ruppersberg, H. Reiter, J. Phys. F: Met. Phys. 12 (1982) 1311– 1325. [53] B. Lohwongwatana, J. Schroers, W.L. Johnson, Phys. Rev. Lett. 96 (2006) 0755031–0755034. ˇ [54] P. Duhaj, P. Svec, Mater. Sci. Eng. A 226–228 (1997) 245–254. [55] R.M. Srivastava, J. Eckert, W. Loeser, B.K. Dhindaw, L. Schultz, Mater. Trans. JIM 43 (2002) 1670–1675. [56] C. Suryanarayana, Rapid solidification, in: R.W. Cahn, P. Haasen, E.J. Kramer (Eds.), Processing of Metals and Alloys, Materials Science & Technology (A Comprehensive Treatment), vol. 15, VCH Publishers, Weinheim, 1991, p. 64. [57] D.V. Louzguine, A. Inoue, J. Mater. Res. 17 (2002) 2112– 2120. [58] M. Kasai, J. Saida, M. Matsushita, T. Osuna, E. Matsubara, A. Inoue, J. Phys. Condens. Matter. 14 (2002) 13867–13877. [59] M. Kasai, E. Matsubara, J. Saida, M. Nakayama, K. Uematsu, T. Zhang, A. Inoue, Mater. Sci. Eng. A 375–377 (2004) 744–748. [60] D.J. Barber, Ultramicroscopy 52 (1993) 101–125. [61] J. Li, X. Gu, T.C. Hufnagel, Microsc. Microanal. 9 (2003) 509–515.