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Ignacio A Figueroa, H. A. Davies, I. Todd
To cite this version:
Ignacio A Figueroa, H. A. Davies, I. Todd. High glass formability for Cu-Hf-Ti alloys with small additions of Y and Si. Philosophical Magazine, Taylor & Francis, 2009, 89 (27), pp.2355-2368.
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High glass formability for Cu-Hf-Ti alloys with small additions of Y and Si
Journal: Philosophical Magazine & Philosophical Magazine Letters Manuscript ID: TPHM-09-Mar-0105.R1
Journal Selection: Philosophical Magazine Date Submitted by the
Author: 22-May-2009
Complete List of Authors: Figueroa, Ignacio A; Univ Sheffield, Dept Mat Engn, Sheffield S1 3JD, S Yorkshire England
Davies, H. A.; University of Sheffield, Department of Engineering Materials
Todd, I.; The University of Sheffield, Engineering Materials Keywords: amorphous metals, copper alloys, metallic glasses, wetting Keywords (user supplied):
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High glass formability for Cu-Hf-Ti alloys with small additions of Y and Si
I. A. Figueroa*, H. A. Davies and I. Todd
Department of Engineering Materials, University of Sheffield, Mappin Street, Sheffield S1 3JD, UK.
*Correspondence Author:
Department of Engineering Materials, University of Sheffield,
Sheffield, S1 3JD, United Kingdom.
Email: i.a.figueroa@sheffield.ac.uk (I.A. Figueroa) Telephone: +44 (0) 114 222 6011
Fax: +44 (0) 114 222 5943 (Received March 2009; final version received )
Abstract
The effects of small substitutions of Si and Y on the glass forming ability of a Cu55Hf25Ti20 glassy alloy are reported and discussed. Fully glassy rods with diameters up to 7 mm and 6.5 mm, were produced for Cu54.5Hf25Ti20Si0.5 and Cu55-xHf25Ti20Y0.3 alloys, respectively. The addition of Si enlarged ∆Tx (=Tx-Tg, where Tg and Tx are crystallisation and glass transition temperatures, respectively) considerably from 25 to 53 K for the Cu54Hf25Ti20Si1 alloy. However, the results showed that the parameters obtained from thermal analysis, such as Trg, ∆Tx and γ[=Tx/(Tg+Tl)] are not reliably correlated with the glass forming ability (GFA), at least for these bulk glass forming alloys. The scavenging effects of the Y and Si, in particular the possibility of Y reducing the oxides, could be responsible for enhancing the GFA. It is proposed that the effectiveness of small additions of Si in enhancing the GFA may be the result of the possible formation of HfSiO4 having a very large negative enthalpy of formation and, as a strong network former, it would form glassy particles which would be ineffective as nucleating agents.
Keywords: amorphous metals; metallic glasses; Cu-based alloys; wetting
1. Introduction
Over the last two decades, bulk glass forming alloys having progressively lower critical cooling rates required for complete amorphisation have been
discovered. This has been largely as a result of extensive and detailed experimental studies, but partly also due to greater understanding of the specific alloying
characteristics that retard crystallisation during cooling from the melt. This has helped
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broaden the range of bulk metallic glass compositions, leading to an increase in research activities on this exciting class of material. Much of this has concentrated on two important issues which need to be addressed in order to promote their widespread usage, namely a) increasing the limiting dimension across which an alloy can be cast into a fully amorphous form and, b) increasing the ductility and mechanical
toughness.
Recently, microalloying has been shown to have substantial effects on the glass formability and the thermal stability for many BMG [1-3]. This involves adding small amounts of alloying additions (usually <2 at.%) to existing bulk glass forming alloys in order to improve their GFA. Our experimental experience indicates that, alloying with > 2 at.% of Al, Mo, Si, B, Nb, V or Y reduced the GFA of Cu55Hf25Ti20 based BMG [4-5]. On the other hand, it was also reported that minor additions (< 2 at.%) of small solute atoms (such as B and Si with atomic radii ∼0.12 nm) or large atoms (such as Y and Sc with radii ∼0.16 nm), in fact enhance the GFA [6-7]. In addition, we have demonstrated that micro additions of Si to Cu-Hf based BMG substantially increase the magnitude of ∆Tx [5]. In the present paper, the effects of additions of ≤ 1 at.% of Si and Y on the GFA and thermal stability of a Cu–Hf–Ti- based alloy are reported and discussed.
2 Experimental Procedure
Cu-based alloy ingots of composition Cu55-xHf25Ti20Mx (M = Si and Y, where x = 0.1, 0.3, 0.5, 0.7 and 1 at.%) were prepared by argon arc melting mixtures of Cu (99.99%
pure), Hf (99.8% pure), Ti (99.98% pure), Si (99.9998% pure) and Y (99.98% pure).
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The alloy compositions represent the nominal values but the weight losses in melting were negligible (<0.1 %). The alloy ingots were inverted on the hearth and re-melted several times to maximise compositional homogeneity. Conical alloy shapes, of length ~50 mm and having a minimum diameter of 1 mm and a maximum of 10 mm, were produced by copper mould suction-casting within the argon arc furnace.
Additionally, ribbon samples of each alloy were produced by chill-block melt
spinning in a sealed He atmosphere at roll speeds in the range 3 m s-1 to 35 m s-1. The structures of the conical and ribbon samples were studied by X-ray diffraction (XRD) (Philips D500 diffractometer with Cu-Kα radiation) and TEM (Philips EM 420 operated at 120 kV). The glass transition and crystallisation temperatures were determined by differential scanning calorimetry (DSC) using a Perkin Elmer DSC-7 at a heating rate of 0.33K s-1. To confirm the reproducibility of the results, at least three samples were measured for each composition. It is worth mentioning that the
criterion used for determining Tg was at the inflection point, i.e.
2 0
2
∂ =
∂ T Cp
, The liquidus (Tl) temperature was determined by differential thermal analysis (DTA) using a Perkin Elmer DTA-7, at a heating rate of 0.33 K s-1. The oxygen levels were
obtained using a LECO analyser and the metallic element concentration of the ternary Cu55Hf25Ti20alloy was determined by the Inductively Coupled Plasma - Optical Emission Spectrometry (ICP-OES) method.
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3 Results
Analysis of the metallic element concentrations, obtained by the Inductively Coupled Plasma technique, showed that the differences between the nominal
(Cu55Hf25Ti20) and actual elemental concentrations (Cu54.978Hf24.952Ti19.94) for one of the alloys were relatively small, showing a maximum difference of ∼0.3 at %. In addition, the weight loss during alloy preparation by electric arc melting, had an average value of 0.04 wt%, which further indicates that the actual elemental
concentrations were very close to the respective nominal values. These results accord with our previous observations for alloys prepared by the same technique.
The Cu55Hf25Ti20 alloy exhibited high glass-forming ability as it is in the compositional vicinity of a deep ternary eutectic; it had a critical glassy diameter, dc, of 4 mm [4-5]. Figure 1 shows the XRD patterns for transverse sections of ingots for the best glass forming compositions for the Cu55-xHf25Ti20Mx (M = Si and Y) alloys.
The diffraction patterns are diffuse, with the absence of distinct crystalline peaks, indicating a fully X-ray glassy phase structure. The microadditions of Si proved to be beneficial for the GFA, as they resulted in an increasing dc to a maximum value of 7 mm at x = 0.5. The addition of x = 0.7 slightly reduced the GFA, to a dc of 6 mm (Figure 2), while at 1 at.% Si the dc decreases to 4 mm. The Y microadditions also enhanced the GFA of the Cu55Hf25Ti20 base alloy. The addition of x = 0.1 increased dc
to 5 mm, as was also observed for Si at the same concentration (Figure 2). The maximum dc was 6.5 mm for Cu54.7Hf25Ti20Y0.3. However, beyond 0.3 at.%Y, dc
decreased substantially down to only 2mm for 1 at.% (Figure 2). Thus, of the
compositions studied, Cu54.5Hf25Ti20Si0.5 (dc = 7 mm) and Cu54.7Hf25Ti20Y0.3 (dc = 6.5 mm) had the highest GFA for the two systems.
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Figures 3a and 3b show the DSC traces for the Si- and Y- microalloyed bulk glassy alloys with diameters of 1.5 mm, together with the data for the Cu55Hf25Ti20 alloy. All the alloys exhibit a distinct glass transition (Tg), followed by the
supercooled liquid region spanning the range ∆Tx before crystallisation at Tx. The values of Tg, ∆Tx and Tx for the glassy Cu55Hf25Ti20 alloy are 713K, 33K and 746K, respectively. It can be seen in the DSC traces that, with the progressively increasing microadditions of Si, ∆Tx increases to ∼55K for the composition Cu54Hf25Ti20Si1, while maintaining the same GFA as the base alloy (4 mm). This implies a high stability of the undercooled liquid state against crystallisation and is encouraging for future development of this glassy alloy as a bulk structural material with the ability to be hot formed to high strains, without devitrification.
There was a marked change in the crystallisation mode from three to two stages between x = 0.7 and x = 1. This suggests that the formation of one of the phases present at lower x content was suppressed.
In contrast to the Si series, the small additions of Y did not result in significant effects on Tg or on the thermal stability of the supercooled liquid region compared with the ternary base alloy. Three-stage crystallisation was observed for all Y contents, although the alloy with x = 0.7 showed only a relatively weak third exothermic peak in the DSC trace at ∼ 860 K, as shown in figure 3b.
The values of ∆Tx are plotted in figure 4a and are also given in table 1. This parameter has been proposed to correlate well with the GFA [8]. However, the data indicate that the correlation between ∆Tx and the critical glassy diameter, dc, in this case was imperfect; for instance, the alloy x = 1 at.%Si having the largest ∆Tx = 53 K
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had a dc of 4mm, whereas the x = 0.5 at. % Si alloy, having a dc = 7 mm, showed a value of ∆Tx = 42 K. As has already been emphasised [5, 9], a large value of ∆Tx
indicates that an alloy can be heated substantially above Tg without inducing
devitrification and thus that the alloy in a highly undercooled liquid state has a high resistance to crystallisation. However, since this temperature range is well below the critical range for glass formation, i. e. the nose of the TTT/CCT curves [9], there is no direct reason why the GFA and the magnitude of ∆Tx should be closely correlated. At higher heating rates (Ṫ), the magnitude of ∆Tx will increase, since Tx is displaced to higher temperatures more rapidly than Tg; this is because crystallisation involves diffusional processes over a longer range than the glass transition. In fact, for Ṫ= 1 K/s, the magnitude of ∆Tx ∼ 67 K. ∆Tx for the Y-containing alloys remained generally constant, with an average value of ∆Tx = 28 ± 2 K, and figure 3 suggests that the crystallisation behaviour did not change significantly over the range 0 - 1 at. %Y.
Figure 5 shows the high temperature DTA curves for samples of the 1.5 mm diameter glassy Cu55-xHf25Ti20Mx (M = Si and Y) alloy rods. Endothermic peak(s) associated with melting are seen for all the alloys and the melting process appears to occur in two stages, indicating slightly off-eutectic compositions in all the cases.
Based on the data shown in figures 3 and 5, the reduced glass transition temperature, Trg (=Tg/Tl) and the gamma parameter, γ=Tx/(Tg+Tl) [10], were calculated for all the Cu55-xHf25Ti20Mx (M = Si and Y) glassy alloys and the data are given in table 1 and shown in figures 4a and 4b, respectively.
Not surprisingly, the small additions of Y and Si, in almost all cases, did not significantly influence the liquidus temperature (Tl) of the base Cu55Hf25Ti20 alloy.
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Only the alloy with 1 at.% Si showed a noticeable change in the melting behaviour, reducing the solidus temperature (Tm) by ∼ 10 K, and increasing Tl by ∼10 K to 1190 K, as shown in figure 5. Thus, the values of Trg for the alloys containing Si were mainly influenced by the increase in Tg with increasing Si. The magnitude of Trg
increased slightly with increasing Si content up to x = 0.7, then more steeply beyond and up to 1 at.%; it has also been shown previously by us to increase for higher Si contents [5]. This contrasts with the behaviour of dc which, although it increased up to x =0.5, then decreased continuously for 0.5 < x ≤ 1.0. On the other hand, the
magnitude of Trg for the alloys with the small additions of Y remained constant, within the experimental error. Similarly, the γ parameter did not show significant correlation with GFA, even though it has been considered to be more useful than ∆Tx
or Trg [10].
4 Discussion
The results in table 1 show that the correlations between the “post mortem”
parameters such as ∆Tx, Trg and γ, obtained from thermal analysis, with the glass forming ability, as measured by the maximum diameter of fully glassy phase dc, are very poor and that, thus, ∆Tx, Trg and γ, are not reliable indicators of GFA for these bulk metallic glasses. The critical cooling rate for glass formation correlates reasonably well with Trg for alloys of low and intermediate GFA, for which
crystallisation kinetics are dominated by homogeneous nucleation [9]. However, the situation for easy glass formers such as the present alloys is more complex and other factors, such as heterogeneous nucleation and the marked variation in the melt viscosity at Tg with alloy composition, tend to preclude the use of Trg alone as a
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controlling parameter [9]. Clearly, in the present case, the thermal analysis and the parameters derived therefrom were unable to explain why such small additions of Si and Y produced strong effects on the GFA. For instance, addition of 0.5 at% of Si to the Cu-Hf-Ti base alloy, in substitution for Cu, resulted in a 75% increase in the critical glassy diameter, up to 7mm.
There is a marked increase in ∆Tx on increasing the Si content of the CuHfTi- based alloy from 0.7% to 1.0 at%. As we reported previously, the supercooled liquid region increases by increasing the Si content, showing a maximum of 75 at 2 at.% Si [5]. This could be interpreted on the basis that the liquid is stabilised by Si in the highly undercooled state and the diffusion rate of Si atoms decreased. According to Lu and Liu [3], small additions of certain atoms, such as Si, generally increase the difficulty of atomic rearrangement, since not only are more atom species required to interdiffuse during heating but also other factors such as the size and electron structure differences play a role. The consequent enlargement of the supercooled liquid region (∆Tx), was also observed by Inoue et al. [11] in the system Zr-Al-Cu with small additions of B. They concluded that the addition of up to 4 at.%B enlarged
∆Tx as a result of the generation of Zr-B pairs with a much stronger nature than the other interatomic pairs and having much longer relaxation time. In the present case, the increase in ∆Tx could be attributed to the formation of Hf-Si pairs which, as can be seen in figure 6, have the strongest bonding of any of the atomic pairs in the system.
The binding energy between free volume and solute atoms may also be a key factor in understanding the impediment to atomic interdiffusion and thus the increased difficulty of atomic rearrangement and consequent enlargement of ∆Tx. The solute
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atoms themselves can be directly influenced by strong binding with the free volume, Thus, Si being the smallest atom, it is probable that it would have the strongest free volume binding energy within a Cu-Hf-Ti-Si alloy. Thus, the Si would bind
preferentially with the free volume, at least up to some saturation level, resulting in reduced diffusivity of Hf and Ti in the alloy and retarding crystallisation to a higher temperature. The supercooled liquid range would therefore be extended to higher temperatures
Figure 6 shows the atomic radii (nm) of the various elements involved in the current study and also the heat of mixing (∆Hm) for the pairs of atomic species. The magnitude of ∆Hm for Si with all of the constituent elements is negative. It has large negative values of ∆Hm with Hf and Ti of -77 and -66 kJ/mol, respectively. The atomic mismatch between Si, Cu and the other elements is also large. The atomic diameters change in the order Hf > Ti > Cu > Si and the atomic size ratios are 1.24 for Hf/Cu, 1.15 for Ti/Cu and 0.90 for Si/Cu. Such mismatches have been proposed as being favourable for easy glass formation [12-13]. Choi-Yim et al. [14], reported that small additions of Si are beneficial to the GFA in alloys similar to those studied here, but containing Zr rather that Hf. It was found that the addition of 0.5–1.0 at% Si increased the maximum thickness of glassy ingots based on Cu–Ti–Zr–Ni from 4 mm to 7 mm. They concluded that Si additions increased the GFA by chemically
passivating impurities that triggered heterogeneous nucleation in the melt. Besides, the relatively small atomic radius of Si could also have occupied some of the
interstitial free spaces and increased the packing density of the liquid structure, which should tend to decrease the free energy and stabilise the liquid phase, while
suppressing the formation of one or more of the crystalline phases [15], as is apparent
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in figure 3a for 1 at.%Si. At high Si contents, there would be a likelihood of all the larger interstitial sites being filled with Si, so that the remaining Si could form compounds, such as silicides [14] which would increase Tl and thus decrease the GFA. However, in the present case, the Si concentration range is considered to be too low to facilitate such an outcome.
In contrast to the Si additions, the Y additions to the Cu-Hf-Ti alloy would be likely to favour the bonding pair Cu-Y in preference to Y–Hf and Y–Ti which have positive heats of mixing. The radius of the Y atom (0.18nm) is the largest of the constituent elements, with the atomic radii being in the order Y>Hf > Ti > Cu; the atomic size ratios are thus 1.24 for Hf/Cu, 1.15 for Ti/Cu and 0.142 for Y/Cu. The Y addition thus broadens the atomic size distribution which would increase the packing density of the alloys. Similarly, the addition of the large Y atoms would tend to retard the diffusion of Cu atoms, which otherwise have a relatively large diffusivity because of their small size. According to Egami [15], the strong atomic bonding between large and small atoms improves the GFA. Nevertheless, it is unlikely that for such small concentrations of Si and Y they would have significant effect on the GFA through changes in packing density and interatomic affinities and an alternative explanation is required.
The negative enthalpies of formation of the various oxides of the constituent metals are given in table 2. The value for SiO2 is intermediate between those for the
‘weak’ oxide formers TiO and Cu2O and the very ‘strong’ oxide former, Ti2O3. Thus, the formation of SiO2 by reduction of HfO2 and Ti2O3 is unlikely. However, there is, in turn, a greater driving force for formation of the ternary oxide HfSiO4 than for
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either HfO2 or Ti2O3, and thus, there is a strong possibility that the Si could combine with the Hf and O to form HfSiO4. This would be expected to be a strong network former, as for other silicates, and thus to have a glassy structure.
It follows from this that, the marked increase in GFA on adding up to 0.5 at.%
Si to the Cu55Hf25Ti20 alloy may result from partial or perhaps, complete removal of Ti2O3 or HfO2 as heterogeneous nucleation sites through the formation of amorphous HfSiO4. The glassy HfSiO4, probably containing also some Cu and Ti network modifiers, would therefore be expected to reduce or nullify the potency of the impurity particle as heterogeneous nucleating agents.
Glassy HfO2 and HfSiO4 have been reported. Normally, they are prepared by the melt quench technique [16]. Recently, hafnia (HfO2) grown on Si substrates has been widely studied as a potential candidate for replacing silica as the gate dielectric in scaled complementary metal oxide semiconductor (CMOS) devices. To avoid crystal orientation, misfit and grain boundary problems at the substrate interface, amorphous thin film substrates are required. The additions of Si increased the crystallisation temperature of the glassy HfO2, leading to the formation of glassy HfSiO4 [17].
The effect of Si would then be to increase the glass forming ability and the initial increase of dc is consistent with this argument. Clearly, the effect saturates beyond an addition of > 0.5 at.% Si, possibly because, at large rod diameters, the cooling rate/ heat transfer rate is not high enough to produce glassy HfSiO4. Another possibility is that the excess of Si could be used to form silicide such as HfSi2 which may act as heterogeneous nucleation sites, leading to a decrease in dc. However, the more likely possibility is that, when all the pre-existing heterogeneous nucleating oxide particles are removed, any additional Si goes into solution and its intrinsic
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effect is then to decrease the GFA. Clearly, further detailed studies are required in order to elucidate the chemistry of the Si addition and its influence on the glass forming kinetics in this alloy.
Table 2 shows the heats of formation of oxides that could be formed with the constituent elements in the Cu-Hf-Ti-Y alloy system. The compound Y2O3 has the largest negative enthalpy of formation; so that it is feasible that Y could reduce any of the oxides that would be present in the absence of Y. Even so, in order to increase the GFA, the Y2O3 would need to be less effective as heterogeneous nucleants that the oxide(s) it replaced. Yttrium addition is well known to be effective for reduction of the oxygen level by surface slag formation during ingot melting [19]. Nagasaka et al.
[19] reported that the presence of a small concentration (0.15 at.%) of Y in the
V91.85Cr4Ti4Y0.15 alloy reduced significantly the oxygen level through the formation of Y2O3 and slagging out onto the surface of the melted ingot. They clearly detected Y2O3 as a skin on the top and sides of the ingot.
For alloys which contain substantial proportions of group IV elements (e.g. Ti, Zr and Hf), which have high affinities for oxygen, significant oxygen pick up is likely, despite our use of a controlled atmosphere. The measured dissolved oxygen contents, determined by LECO analysis, for some of the alloys investigated, i.e.
Cu55Hf25Ti20, Cu54.5Hf25Ti20Si0.5 and Cu54.5Hf25Ti20Y0.5, were 316, 252 and 233 ppm by weight, respectively. These results show that the addition of Si and Y reduced the oxygen level by 54 ppm for the Si- and 83 ppm for the Y- containing alloys
(equivalent to ~17 and ~ 26 %, respectively). This indicates that oxides of the minor elements are being formed.
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The formation of small oxide particles within the melt can be considered to be detrimental to the glass forming ability, as they could act as heterogeneous nucleation sites. The fact that the glass forming abilities of Pd-based alloys have been shown to be improved by fluxing with B2O3 indicates that such a heterogeneous nucleation mechanism is normally operative in glass forming alloys [20]. Drehman and Greer [20] and, later, Nishiyama and Inoue [21] found that, depending on the processing conditions, crystallisation in Pd40Ni40P20 alloy and in a PdCuNiP alloy, respectively, can be either due to growth from internal nuclei or from surface impurities. It was concluded that heterogeneous nucleation was a substantial limitation on the GFA in the Pd-Ni-P alloys [20].
In the case of Y-containing alloys, as indicated earlier in the discussion, the likelihood is that the Y would have removed the pre-existing oxides or formed in preference to the other oxides. These may have floated off to the ingot surface [19].
However, if not, their effect as particulate in the molten alloy on the crystal nucleation behaviour needs to be considered. Naidich et al. [22] showed that the wetting angles of liquid Cu, Ag and Au on rare earth oxides are in the range ∼ 120 – 140o on Y2O3
(figure 7a). However, for Cu-Ti alloys (figure 7b), the wetting angle on Y2O3 is rapidly reduced with increasing Ti concentration, for Cu-20at.% Ti the wetting angle is reduced to ∼ 70 – 80o.
The effect of the angle of contact of heterogeneous nucleants has been considered by Lewis and Davies [23] and, more recently, by Nishiyama and Inoue [21]. Lewis and Davies considered for a model glass forming alloy and assuming a spherical cap model, the influence of the angle of contact, θ, on the critical cooling rate for glass formation, Rc (Figure 8a). It can be seen that, on lowering θ below a
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critical value of ∼80° for low and intermediate values of Trg ≲ 0.5, the effect of the
heterogeneous nucleants is relatively small since the homogeneous nucleants have a dominating influence. In contrast, as Trg is increased to a typical value for bulk glass formers, below the critical value of θ, the heterogeneous nucleation becomes very dominant, with Rc increasing by ∼ 4 orders of magnitude as θ → 0. This adds further weight to the unreliability of Trg as a figure of merit for bulk glass forming alloys on account of the likely variability in the concentration and potency of heterogeneous nucleants from alloy to alloy. It should be pointed out that this model for the effects of heterogeneous nucleation needs to be reassessed for bulk glass formers, as the
viscosity at the liquidus temperature was assumed, for the purpose of comparison, to be independent of Trg and, for this reason, the values of cooling rates calculated Rc, calculated on the basis of this model are excessively high for bulk glass forming alloys in comparison with experimentally estimated Rc.
No experimental data were found for the angle(s) of contact between crystalline phase(s) and the HfO2 or Ti2O3 in Cu-Hf-Ti melts. We speculate that Hf and Ti oxides are more effective heterogeneous nucleating agents for Cu alloys than their replacement Y2O3; thus, the Rc would be expected to decrease on addition of Y, depending on the concentration of nuclei and the precise value of θ. This would explain the enhancement of GFA on initially adding Y to the Cu-Hf-Ti alloy until all the Hf and Ti oxides are transformed and/or avoided, beyond which the effect of Y is evidently to reduce the GFA when present as a solute element in the alloy.
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Nishiyama and Inoue [21] reported that the crystallisation of the
Pd40Cu30Ni10P20 alloy is dominated by a heterogeneous nucleation mechanism and the wetting angle of these nucleants (Figure 8b). The foreign substrate that induced nucleation of crystalline embryos in the undercooled melt were predicted to have a critical wetting angle, θ, ∼63° for this particular system. Therefore, the total
nucleation frequencyIvtotal increases with decreasing θ. Furthermore, the computations indicate that the peak of Ivtotal is rapidly elevated, broadened and shifted to higher temperature at low values of θ. If θ = 0°, representing perfect wetting, the molten alloy could not be undercooled below Tm, even at high cooling rates.
Thus, to account for the observation of GFA going through a maximum and falling beyond 0.3 at.%, it would appear necessary to invoke a saturation effect of the Y, whereby all the heterogeneous nucleants are transformed to Y2O3 and, thereafter, the excess of Y in solution having a detrimental intrinsic effect on the GFA. However, the manner in which heterogeneous nucleation influences crystallisation and therefore limits the GFA is not yet fully understood as there are many uncertainties involved. It is clear that further analysis of nucleation and growth in this particular alloy with dopant additions is required.
5 Conclusions
The small substitutions of Si and Y for Cu in the ternary Cu55Hf25Ti20 alloy increased considerably the GFA, as fully glassy rods with diameters up to 7 mm and 6.5 mm, were produced for Cu54.5Hf25Ti20Si0.5 and Cu55-xHf25Ti20Y0.3 alloys,
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respectively. The addition of Si enlarged the ∆Tx considerably from 25 to 53 K for the Cu54Hf25Ti20Si1 alloy. The enlargement of ∆Tx suggests that the liquid is stabilised by Si, at least in the highly undercooled state, and could be attributed to the formation of Hf-Si pairs which have the strongest bonding in the system. Nevertheless, the results show that the parameters obtained from thermal analysis, such as ∆Tx, Trg and γ generally do not correlate with GFA, at least for the lightly doped bulk glass forming alloys. The scavenging effect of the Y and Si, in particular the possibility of Y reducing the oxides, could be responsible for enhancing the GFA. It is proposed that the effectiveness of small additions of Si in enhancing the GFA may be related to the formation of HfSiO4 which has a large negative heat of mixing. This phase would be expected to be a strong network former, as for other silicates, and thus to solidify to a glassy structure prior to the quench. It is probable that the HfSiO4 contains some Cu and Ti as network modifiers, such that both the potency of the original TiO2 and HfO2
heterogeneous nucleants is reduced or nullified by their replacement by non- crystalline HfSiO4.
Acknowledgements
IAF is grateful for the financial support of the Mexican Council of Science and Technology (Conacyt), through PhD Research Scholarship No. 205146. Valuable technical support provided by Mr. P.
Hawksworth is also acknowledged.
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References
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496 .
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[8] A. Inoue, “Bulk Amorphous Alloys; Preparation and Fundamental
Characteristics”, TransTech Publications-Switzerland, 1998, Mater. Sci. Foundations No. 6.
[9] H.A. Davies, Metallic Glass Formation Revisited, Proceedings of the 4th International Workshop on Non-Crystalline Solids, World Scientific Press, Madrid-Spain,1994, p. 3.
[10] Z.P. Lu, C.T. Liu, Phys. Rev. Lett., 91 (2003) p. 115505.
[11] A. Inoue, T. Negish, H. Kimura and T. Aoki, Mater. Trans. JIM. 38 (1997) p.185.
[12] A. Inoue, Acta Mater. 48 (2000) p. 279.
[13] B. C. Giessen, Glass formation diagrams: A two parameter representation of readily glass forming binary alloys systems, Proc. 4th Int. Conf. on Rapidly Quenched Metals, Senday, Japan, 1981, p. 213.
[14] H. Choi-Yim, R. Busch, W.L. Johnson, J. Appl. Phys. 83-12 (1998) p. 7993.
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[18] O. Kubaschewski and C. B. Alcock, Metallurgical Thermochemistry, 5th Ed., Pergamon press, 1979.
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Nucl. Mater. 367–370 (2007) p. 823.
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[21] N. Nishiyama, and A. Inoue, Acta Mater. 47 (1996) p. 1487.
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Table 1. Critical glass forming diameter, dc, supercooled liquid region (∆Tx), reduced glass transition temperature (Trg) and gamma parameter (γ) for the Cu55-xHf25Ti20Mx
(M = Si and Y) glassy 1.5 mm conical rod diameter.
Alloy dc mm
Tg K
Tx K
Tl K
∆∆
∆∆Tx
K
Trg γγγγ Si
0.0 4 725 750 1190 25 0.609 0.392
0.1 5 727 760 1190 33 0.611 0.396
0.3 6 731 765 1198 34 0.610 0.397
0.5 7 736 770 1195 34 0.616 0.399
0.7 6 736 773 1190 37 0.618 0.401
1.0 4 745 798 1205 53 0.618 0.409
Y
0.1 5 725 750 1190 25 0.609 0.392
0.3 6.5 725 750 1190 25 0.609 0.392
0.5 5.5 725 750 1187 25 0.611 0.392
0.7 3 725 750 1185 25 0.612 0.393
1.0 2 725 755 1185 30 0.612 0.395
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Table 2. Heat of formation of various oxides of the constituent metals [18]
Substance Enthalpy of formation (KJ/mol)
CuO -155.0
Cu2O -167.0
TiO -543.0
SiO2 -911.0
TiO2 -945.4
HfO2 -1114.0
Ti2O3 -1521.0
Y2O3 -1907.0
HfSiO2 -2010.0
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Figure captions
Figure 1. XRD patterns of the best glass forming compositions obtained in the Cu55- xHf25Ti20Mx (M = Si and Y, where x = 0.1, 0.3, 0.5, 0.7 and 1 at.%) alloys.
Figure 2. Glass forming section thickness, dc, for the Cu55-xHf25Ti20Six and Cu55- xHf25Ti20Yx (x = 0.1, 0.3, 0.5, 0.7 and 1 at.%)
Figure 3. DSC curves cast a) Cu55-xHf25Ti20Six and b) Cu55-xHf25Ti20Yx (x = 0.1, 0.3, 0.5, 0.7 and 1 at.%) glassy 1.5 mm diameter rods, at a heating rate of 0.33 K s-1. Figure 4. a) Supercooled liquid region (∆Tx) and reduced glass transition temperature (Trg) and b) Gamma parameter (γ), as a function of Solute “M” content for the Cu55-
xHf25Ti20Mx (M = Si and Y ) glassy 1.5 mm diameter rods.
Figure 5 DTA curves of cast a) Cu55-xHf25Ti20Six and b) Cu55-xHf25Ti20Yx (x = 0.1, 0.3, 0.5, 0.7 and 1 at.%) glassy 1.5 mm diameter rods, at a heating rate of 0.33 K s-1. Figure 6 Relationship of heat of mixing and atomic radii among constituent elements in the Cu–Hf–Ti –(Si,Y) alloy system.
Figure 7. a) Angles of wetting of RE oxides and aluminium oxide by melts of metals:
(1) Cu; (2) Ag; (3) Au; (4) Sn; (5) Ge; T = 1150°C and b) Concentration dependences of wetting by Cu-Ti melts of oxides (1) Al2O3, (2) Y2O3, (3) Er2O [22].
Figure 8. a) Theoretical variation of cooling rate (Rc) with contact angle (θ) for various Trg [23], and b) Total nucleation frequency Ivtotal as a function of temperature based on the assumptions of various wetting angles (θ) between 0° and 180° [21].
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Figure 1
294x177mm (96 x 96 DPI)
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Figure 2
208x129mm (96 x 96 DPI)
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Figure 3a
191x209mm (96 x 96 DPI)
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Figure 3b
201x243mm (96 x 96 DPI)
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Figure 4 a
259x125mm (96 x 96 DPI)
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