Journal of Non-Crystalline Solids 354 (2008) 3659–3665

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Journal of Non-Crystalline Solids

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A new composition zone of bulk metallic glass formation in the Cu–Zr–Titernary system and its correlation with the eutectic reaction

Chun-Li Dai a, Hua Guo a, Yi Li b, Jian Xu a,*

a Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, Chinab Department of Materials Science and Engineering, National University of Singapore, Engineering Drive 1, Singapore 117675, Singapore

a r t i c l e i n f o

Article history:Received 31 October 2007Received in revised form 9 April 2008

PACS:61.43.Dq61.25.Mv64.70.pe

Keywords:AlloysTransition metalsGlass transitionScanning electron microscopyPhases and equilibriaCalorimetryGlass transitionX-ray diffraction

0022-3093/$ - see front matter � 2008 Elsevier B.V. Adoi:10.1016/j.jnoncrysol.2008.04.013

* Corresponding author. Tel.: +86 24 23971950; faxE-mail address: jianxu@imr.ac.cn (J. Xu).

a b s t r a c t

A new composition region of bulk metallic glass formation, around Cu52Zr40Ti8, was discovered in the Cu–Zr–Ti ternary system, for which monolithic bulk metallic glass rods of 4 mm in diameter can be fabricatedusing copper mold casting. The solidification of the Cu52Zr40Ti8 deeply-undercooled liquid mainly under-goes a univariant eutectic reaction, (L ? Cu10Zr7 + CuZr), even though this composition was predicted tobe a ternary eutectic point (L ? Cu10Zr7 + CuZr + Cu2ZrTi) by CALPHAD calculations. With respect to thedeep-eutectic reaction of (L ? Cu10Zr7 + CuZr) in the Cu–Zr binary alloys, alloying of Ti has a significanteffect on further stabilizing the liquid, as indicated as a drop of the univariant eutectic groove, limitingthe coupled growth of two crystalline phases, hence increasing the glass-forming ability.

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1. Introduction

Among the family of bulk metallic glasses (BMGs), Cu-based al-loys are particularly interesting for practical application as newstructural materials due to their significant advantages, includingthe low-cost, high fracture strength around 2 GPa often coexistingwith visible ductility, and the feasibility for formation of BMG-based composites [1–10]. So far, Cu-based BMGs with a critical size(Dc) of centimeter-scale have been discovered in several alloy sys-tems, such as the quaternary Cu46Zr42Al7Y5 [7], Cu43Zr43Al7Be7 [8],Cu44.25Ag14.75Zr36Ti5 [9], Cu45Zr42.55Al9Y3.45 (Dc = 14 mm) [10] andeven ternary Cu49Hf42Al9 [11].

The first Cu-based BMG, Cu47Ni8Ti34Zr11 (Vit 101) with a Dc of4 mm prepared by copper mold casting, was reported by Lin andJohnson [1]. They started from a Cu55Zr10Ti35 composition closeto a ternary invariant eutectic point, shown as E3 in Fig. 1. The pro-jection of liquidus surface of this Cu–Zr–Ti ternary phase diagramwas experimentally established by Woychik and Massalski [12],and actually shows three ternary invariant eutectic points denoted

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as Ei (i = 1,2,3, respectively), marked as red filled squares. Then, Nipartial substitution for Cu in the Cu55Zr10Ti35 (Dc = 0.5 mm only)leads to a further increase in the glass-forming ability (GFA) untilthe Dc is up to 4 mm. Later, Inoue and his collaborators discovereda BMG (Dc = 4 mm), with quite different composition Cu60Zr30Ti10

in the same ternary Cu–Zr–Ti system [2]. However, no consider-ation for its correlation with eutectic reaction was addressed intheir work. Recently, the phase diagram of Cu–Zr–Ti ternary wasrevisited using the CALPHAD approach [13]. From the calculation,it was claimed that there exist five ternary invariant eutectic pointsin this system, herein denoted as E0i (i = 1,2,3,4,5, respectively) andmarked as blue filled circles in Fig. 1 for comparison. Among theseeutectic points given by CALPHAD, it is of interest to note that theE03 point at the Cu52Zr40Ti8 composition is for a ternary eutectic,(L ? Cu10Zr7 + CuZr + Cu2ZrTi), and the eutectic temperature wasestimated to be �1112 K. However, this prediction has not beenexperimentally verified yet. Moreover, it is of interest to learnwhether or not the alloys surrounding this ‘‘eutectic composition’’(which is quite different from the above two composition zones inthis ternary system) potentially have a high GFA owing to theirdeep-eutectic feature. It is worthwhile to emphasize that for a gi-ven ternary system, and even for a simple binary one, it is probable

Fig. 2. DSC scans during the heating and cooling associated with the melting andsolidifying events, respectively, for the Cu52Zr40Ti8 (E03) arc-melted ingot. Insetdisplays a curve fitting result for the curve during heating.

Fig. 3. XRD patterns of arc-melted ingots for the Cu52Zr40Ti8 (E03) together with twoCu–Zr binary alloys, eutectic Cu56Zr44 (e3) and hypereutectic Cu52Zr48 with the sameCu content as Cu52Zr40Ti8 but without Ti.

Fig. 1. Projection of liquidus surface of Cu–Zr–Ti ternary phase diagram establishedby Woychik et al. [12], where the ternary invariant eutectic points (Ei, i = 1,2,3) aremarked as red filled squares. Ternary invariant eutectic points (E0i , i = 1,2,3,4,5)given by CALPHAD calculation (Arroyave et al. [13]) are marked as blue filled circles.Alloys involved in this work situate in the orange triangle region. (For interpretationof the references to color in this figure legend, the reader is referred to the webversion of this article.)

3660 C.-L. Dai et al. / Journal of Non-Crystalline Solids 354 (2008) 3659–3665

that there exist several BMG-forming composition zones withindifferent crystalline phase boundaries depending on differenteutectics [14]. It is expected that the GFA inside each zone is di-rectly correlated with the nature of the eutectic reaction therein.

In this paper, it will be experimentally tested whether theCu52Zr40Ti8 composition (E03) is a ternary invariant eutectic point,primarily based on the melting behavior of the alloy. This will befollowed by a study on the phase selection in the sample solidifiedat a cooling rate slightly lower than the critical cooling rate forBMG formation. Furthermore, the GFA and the feasibility of BMGformation at this composition, together with the compositiondependence of the GFA in the vicinity of this alloy, will beinvestigated.

2. Experimental

Commercial elemental pieces with purities of 99.9 wt% or betterwere used as starting materials. The master alloy ingots of 20 g inweight with the nominal composition (in atomic percentage) wereprepared by arc-melting under a Ti-gettered argon atmosphere ina water-cooled copper crucible. The alloy ingots were melted severaltimes to ensure compositional homogeneity. The weight change ofthe master alloys before and after arc-melting is less than 0.1 wt%.The ribbon samples were prepared in an argon atmosphere by induc-tion melting the master alloy in a quartz tube and ejected it onto asingle-roller using a melt spinner. The surface velocity of the copperroller was of 39 m s�1. The as-quenched ribbons were approximately45 lm thick and 2 mm wide. Smaller bulk samples (diame-ter6 2 mm) are fabricated using suction casting in a mini-arc mel-ter. For the bulk samples with a diameter between 3 and 5 mm,the master alloy was re-melted in a quartz tube using inductionmelting and injected in a purified inert atmosphere into the coppermold that has internal rod-shaped cavities of about 50 mm in length.

The arc-melted ingots and as-cast rods were sectioned trans-versely and polished for scanning electron microscopy (SEM)observation and X-ray diffraction (XRD) analysis. SEM observationof the samples was carried out in a LEO Supra 35 scanning electronmicroscope. The local compositions were semi-quantitatively

determined using an energy dispersive X-ray spectrometer (EDX)attached to the SEM. XRD analysis was performed with a RigakuD/max 2400 diffractometer with monochromated Cu Ka radiation.

The glass transition and crystallization behavior of the glassyribbons or rods were investigated in a Perkin-Elmer differentialscanning calorimeter (DSC-diamond) with an alumina containerunder flowing purified argon at a heating rate of 0.67 K s�1. A sec-ond run under identical conditions was used to determine thebaseline after each run. To confirm the reproducibility of the exper-imental results, at least three samples have been measured foreach composition. All the measurements of the glass transitiontemperature (Tg) and onset temperature of crystallization events(Tx1) were reproducible within the error of ±1 K. The heat of crys-tallization DHx for the glassy phase was determined by integratingthe area under the DSC curve. The melting behavior of the alloyswas measured in a Netzsch 404 DSC with an alumina container,using the heating and cooling rates of 0.33 K s�1.

3. Results

3.1. Certification of the Cu52Zr40Ti8 (E03) to be a ternary eutectic point

Fig. 2 shows the DSC curves for the arc-melted Cu52Zr40Ti8 (E03)alloy near and above its melting temperature during heating and

C.-L. Dai et al. / Journal of Non-Crystalline Solids 354 (2008) 3659–3665 3661

cooling. Both melting process and solidification of the alloy wereperformed mainly through three steps rather than a single event,as seen an inset in Fig. 2. The curve for the melting event is fittedwith EM Gauss function using an Origin peak-fitting program. Thisimplies that this composition is not really located at a ternaryinvariant eutectic point as predicted by CALPHAD approach inRef. [13]. The onset temperature of the melting (Tm) and liquidustemperature (TL) for the alloy was determined to be 1108 and1168 K, respectively. It was noted that the Tm is about 55 K lowerthan the eutectic temperature (1163 K) of the binary eutectic reac-tion of (L ? Cu10Zr7 + CuZr) (denoted as e3 hereafter) in the Cu–Zrbinary subsystem (see phase diagram in Ref. [15]).

It was noticed that the cooling rate for the arc-melted ingot isslightly lower than that the BMG formation required under con-ventional copper mold casting. For the Cu–Zr–Ti alloy, it is esti-

Fig. 4. Backscattered SEM images of arc-melted ingots for the (a) Cu52Zr40Ti8, (b)binary eutectic Cu56Zr44 (e3) and (c) binary hypereutectic Cu52Zr48.

mated to be on the order of K s�1 [16]. Therefore, solidifiedmicrostructure of the arc-melted ingot directly reflects the phaseselection during crystallization of the alloy melt, which competeswith the glass formation. In this sense, the eutectic reaction thatthe alloy melt undergoes during cooling can be revealed. Fig. 3illustrates the XRD pattern of the arc-melted Cu52Zr40Ti8 alloy.For comparison, also plotted in the figure are the XRD patterns oftwo selected binary alloys without Ti, which are Cu56Zr44 at eutec-tic point e3 and a hypereutectic Cu52Zr48 (corresponding toCu52Zr40Ti8 if the latter is treated as the Cu–(Zr,Ti) pseudo binary),prepared under identical conditions. For the Cu52Zr40Ti8 ternary,crystalline phases Cu10Zr7 (orthorhombic) and CuZr can be identi-fied. The crystalline CuZr is present in two forms with differentstructures, the B2 structure (CsCl-type, bcc based) with a higherfraction, and the B190 basic martensite structure (NiTi-martens-ite-type) [17] with a lower fraction. It is noteworthy that the ter-nary compound Cu2ZrTi (MgZn2-type Laves phase) is notdetectable. In the cases of two Cu–Zr binary alloys, the groundstate phase Cu10Zr7 and metastable phase CuZr with B2 structureand basic martensite structure coexist. It indicates that the phaseselection in the ternary alloy is nearly the same as in the two bin-ary alloys.

Figs. 4(a)–(c) show the backscattered SEM images of the arc-melted ternary Cu52Zr40Ti8 and the binary Cu56Zr44 and Cu52Zr48 al-loys. The lateral surface of the sample was polished for SEM obser-vation, under which it was shown that the microstructure isuniform throughout the sample from the top to the bottom. Twophases with different contrasts can be seen in Fig. 4(a) for the ter-nary alloy, marked as A and B, respectively. Such a microstructureis a typical eutectic morphology, i.e. a mixture of two non-facettedrod-like grains resulting from the co-precipitation and coupledgrowth during solidification. By using the EDX analysis, chemicalcomposition for the area of A and B was determined to be Cu58Zr36-Ti6 and Cu49Zr44Ti7, respectively. Then, the phase marked as the Aand B can be identified as the Cu10Zr7 and CuZr phase, respectively,each containing a small amount of Ti. In other words, the Ti ofabout 6 and 7 at.% can be dissolved in the Cu10Zr7 and CuZr phase,respectively. It means that the Cu10Zr7 and CuZr phase in the ter-nary alloy are no longer to be a stoichiometric compound likethe case of Cu–Zr binary. The volume fraction of Cu10Zr7 phase isslightly higher than that of CuZr phase. For the binary eutecticCu56Zr44, CuZr dendrites (lighter area) are dispersed in the Cu10Zr7

matrix, as shown in Fig. 4(b) (marked as B and A, respectively),with a size finer than the case of the ternary. As seen in Fig. 4(c),a typical hypereutectic morphology is observed in the Cu52Zr48,where the coarser CuZr phase (lighter area) precipitated duringsolidification as the primary phase and reached a high volume frac-tion (�80 vol.%). Compared with the E03 alloy, apparently the 8% Tisubstitution for Zr in the binary drives the solidified microstructurein the arc-melted sample from hypereutectic towards two-phasecoexisting eutectic morphology, under the current cooling rate.

3.2. Glass-forming ability of the alloys around Cu52Zr40Ti8 (E03)

In terms of the eutectic-like feature of the Cu52Zr40Ti8 (E03), itsGFA were examined by rapidly quenching of the melt. The XRDpattern of the melt-spun ribbon of this alloy exhibits typical amor-phous nature without any detectable crystalline phase (not shownhere). Additional evidence is that the DSC curve of the ribbon (notshown here) shows a clear endothermic signal associated withglass transition and several sharp exothermic peaks resulting fromcrystallization. From the curve, the Tg and Tx1 of the glassy alloy aredetermined to be 695 and 726 K, respectively. The width of thesupercooled liquid region DTx (DTx = Tx1 � Tg) and reduced glasstransition temperature Trg (Trg = Tg/TL) are then 31 K and 0.59,respectively. Such large values of the DTx and Trg imply a feasibility

Fig. 5. Composition map of BMG formation for as-cast rods in the case of 4-mm-diameter around the Cu52Zr40Ti8 (E03). The full and half open symbols represent thefully and partially glassy phases, respectively. The assumed tie-lines between thecompounds are drawn as the dash lines.

Fig. 6. (a) XRD patterns taken from the cross-sectional surface and (b) DSC scans (ata heating rate of 0.67 K s�1) of the as-cast 4-mm-diameter rods for the Cu–Zr–TiBMG-forming alloys.

Table 1Thermal properties of Cu–Zr–Ti BMGs (Dc = 4 mm) fabricated using copper moldcasting, determined from DSC measurements

Alloys (at.%) Tg (K) Tx1 (K) DTx (K) DHx (kJ/mol) Tm (K) TL (K) Trg

Cu52Zr42Ti6 698 745 47 6.32 1134 1193 0.58Cu51Zr42Ti7 704 738 34 6.25 1128 1186 0.59Cu51Zr41Ti8 700 727 27 6.18 1105 1180 0.59Cu52Zr40Ti8 (E03) 695 726 31 6.04 1108 1168 0.59Cu51Zr40Ti9 696 724 28 6.11 1111 1166 0.59

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to fabricate the alloy into the BMG. Indeed, the Dc of BMG forma-tion for this alloy can reach to 4 mm by virtue of casting the rodsamples in several diameters up to 5 mm.

Considering that the GFA of alloys in a given system is stronglycomposition-dependent even in the case of 1 at.% change [18–20],we used the new BMG-former Cu52Zr40Ti8 as a starting point to lo-cate the best BMG-forming composition in this vicinity. The com-position-dependent Dc under copper mold casting was mappedout by varying the composition with an interval of 1 at.%.

Fig. 5 displays the BMG-forming composition map for 4-mm-diameter as-cast rods. Within the selected triangular region, threecompounds (phases) exist, CuZr, Cu10Zr7 and Cu2ZrTi. The assumedtie-lines between the compounds are drawn as the dash lines. Thelocation of the investigated range in the Cu–Zr–Ti ternary system ismarked as a small triangle in Fig. 1. In the map, the half open andfull symbols represent partially glassy and fully glassy phases,respectively. As seen in Fig. 5, fully glassy rods with 4-mm-diame-ters are found at five composition including E03. But, all of themcompletely crystallize when the diameter is increased up to5 mm (not shown here). It is noticed that the GFA of the alloys ismore sensitive to the change of the ratio between Cu and Zr thanTi concentration.

Fig. 6(a) displays the XRD patterns taken from the cross-sec-tional surfaces of 4-mm-diameter as-cast rods for the Cu52Zr42Ti6,Cu51Zr42Ti7, Cu51Zr41Ti8, Cu52Zr40Ti8 (E03) and Cu51Zr40Ti9 alloys.The typical diffusive maxima reflecting the amorphous feature,without detectable crystalline peaks within the XRD resolution,demonstrate the complete glass formation at this size. The corre-sponding DSC traces for these glassy rods are given in Fig. 6(b).In all cases, a pronounced endothermic signal resulting from theglass transition and multiple-step crystallization events are ob-served. The thermal properties measured from DSC curves for thefive alloys are summarized in Table 1, including the Tg, Tx1, DTx

and DHx together with the Tm, TL and Trg (see below). For theseglasses, the Tg is nearly the same but the Tx1 is composition-depen-dent, such that the DTx varies within a range of 27–47 K and the Trg

remains around 0.6, even though the Dc of the alloys remains at thesame level.

For the purpose of comparison, the GFA of Cu–Zr binary alloyswithin the composition range of Cu58.82Zr41.18–Cu50Zr50, only re-lated to the e3 reaction, is investigated as well, with a compositioninterval of 1 at.%. It is found that within a wide region betweenCu56Zr44 (the e3 point) and Cu50Zr50, 1-mm-diameter glassy rodscan be formed. Fig. 7 shows the XRD patterns of the 1.5-mm-diam-eter as-cast rods of Cu51Zr49 together with two neighboring alloys,Cu52Zr48 and Cu50Zr50, indicating that the optimal BMG-formingcomposition is located at Cu51Zr49 (Dc � 1.5 mm), at which only asmall amount of crystalline phases remains. An inset in Fig. 7displays the DSC curves of the melt-spun ribbon and 1.5-mm-diameter rod for the Cu51Zr49 alloy. The Tg, Tx1, DTx and DHx forthe 1.5-mm-diameter rod was determined to be 680, 730 and50 K and 6.39 kJ/mol, respectively. The nearly identical thermalproperties of the samples with different sizes up to 1.5 mm indi-cate that the Dc for this composition is almost 1.5 mm. Apparently,it is this binary eutectic reaction that lays the foundation of

Fig. 7. XRD patterns taken from the cross-sectional surface of the as-cast 1.5-mm-diameter rods for the binary Cu50Zr50, Cu51Zr49 and Cu52Zr48 alloys and a insetshowing the DSC scans (at a heating rate of 0.67 K s�1) for the melt-spun ribbon and1.5-mm-diameter rod for Cu51Zr49 alloy.

Fig. 8. (a) DSC scans of near and above melting temperatures during heating (at aheating rate of 0.33 K s�1) and cooling for the Cu93�xZrxTi7 (x = 37,39,41,43) seriesalloys as the typical. (b) The liquidus and solidus surfaces established using the dataof TL and Tm from DSC measurements for the 25 relevant alloys. The compositionswith Dc = 4 mm of BMG formation are marked on the projection plane forcomparison.

C.-L. Dai et al. / Journal of Non-Crystalline Solids 354 (2008) 3659–3665 3663

enhancing the BMG formation by further Ti addition, resulting inthat the Dc increased from �1.5 mm for the binary up to 4 mmfor the ternary.

3.3. Liquidus and solidus of the BMG-forming alloys

Fig. 8(a) shows the DSC curves of several representative alloys,Cu93�xZrxTi7 (x = 37,39,41,43), near and above their respectivemelting temperatures, during heating and cooling. The Tm and TL

are marked with arrows. Using similar data from the 25 relevantcompositions with an interval of 1 at.%, a schematic diagram ofthe trend of the Tm and TL as a function of composition is plottedin Fig. 8(b). For the alloys with a Ti content less than or equal to7 at.%, only two endothermic peaks associated with melting ap-peared in the DSC scans, implying that the ternary invariant eutec-tic reaction is not involved. Exceeding 7 at.% for the Ti content, themelting event occurs over three steps, similar to that shown inFig. 2. As shown in Fig. 8(b), the Tm of the alloys only changesslightly when the Ti content is fixed, but a pronounced depressionof Tm occurs as the Ti content was increased, dropping from 1136 Kat 5%Ti to 1110 K at 9%Ti. With respect to the binary invariant eu-tectic temperature (e3, 1163 K), the Tm value dropped about 27–53 K. Evidently, introducing Ti significantly stabilizes the liquid rel-ative to the Ti-free binary Cu–Zr alloys. Similarly, it is shown thatthe TL also decreased as the Ti content was increased. In addition,it is noticed that the TL variation shows a ‘U’-shape dependenceon the Cu/Zr fraction. The locus of the minimum TL values exhibitsa trend like eutectic valley (see Fig. 8(b)). It is shown that the TL

drops from 1179 K at the Cu55Zr39Ti6 down to 1158 K at theCu53Zr38Ti9. These findings demonstrate that within the investi-gated composition region, solidification of the liquids upon coolingmainly undergoes a univariant eutectic reaction of (L ? Cu10Zr7 +CuZr).

In addition, five compositions with Dc = 4 mm are marked asellipses on a projection plane in Fig. 8(b). Note that these alloysare not just located inside the univariant eutectic valley, but skewto the side with the higher TL (CuZr phase terminal). Such a behav-ior has also been observed before in other BMG-forming systems[19,20], and is understood based on the skewed eutectic coupledzone about the eutectic point, due to the competitive growth ofthe participating phases [21]. Similarly, this explanation can be ex-tended to the Cu–Zr binary alloys related to the e3 reaction, forwhich the best BMG-former is located at Cu51Zr49 rather than atthe eutectic point Cu56Zr44.

4. Discussion

As seen in Fig. 2, the melting and solidifying for the Cu52Zr40Ti8

alloy at a rate of 0.33 K s�1 proceed through three steps. Solidifica-tion pathway of the alloy melt is expected to undergo the followingsequence: primary crystallization to solidifying CuZr phase, thenthe univariant eutectic of (L ? Cu10Zr7 + CuZr), and finally a ter-nary invariant eutectic of (L ? Cu10Zr7 + CuZr + Cu2ZrTi). A largerfraction of melting latent heats originating from the second stepindicates that the univariant eutectic reaction is the dominantevent with respect to the others. Furthermore, it was demonstratedthat this composition is actually far away from the true ternaryinvariant eutectic point.

As shown in the microstructure of the Cu52Zr40Ti8 arc-meltedalloy, formation of Cu10Zr7 and CuZr two-phase eutectic indicatesthat both the CuZr primary crystallization and the ternary eutecticreaction can be suppressed under this non-equilibrium condition.Additionally, it was found in the Al–Cu–Ag system that solidifica-tion of ternary alloys with a composition close to the univarianteutectic reaction under certain conditions leads to two-phase pla-nar growth, showing a morphology similar to what is known frombinary (invariant) eutectic growth [22,23]. As suggested by Boett-inger [21], microstructures may in many cases be the result ofcompetition between various forms of crystal growth, and the

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formation of metallic glass can be promoted the limitations on eu-tectic growth kinetically controlled by diffusion. Consequently, ourfindings that the alloys along the univariant eutectic groove favorthe glass formation can be understood in terms of the limitationof (Cu10Zr7 + CuZr) two-phase coupled growth during undercoolingof the melt.

Recently, the formation of BMGs (Dc = 1–2 mm) was reported atseveral compositions in the Cu–Zr binary system [7,8,19,24,25]. Asshown in the phase diagram [15], the system contains five eutecticreactions involving different crystalline phases. For comparison,the current BMG-forming compositions and corresponding eutec-tic reactions are summarized in Table 2. Re-examination of thissystem confirms that for the alloys related to the e3 reaction(Cu50Zr50–Cu58.82Zr41.18), the optimal composition for glass forma-tion is off-eutectic, located at Cu51Zr49. Note that it is only 1 at.% faraway from equiatomic composition Cu50Zr50, where the liquid isjust above a congruent melting compound CuZr. The compositionCu50Zr50 was found to be a marginal BMG which appears ‘X-rayamorphous’ (no detectable Bragg peaks) when cast to up to 1-mm-diameter, but shows a fine dispersion of nanocrystals withdimensions of the order of 5 nm under TEM [6,26]. Thereby, theGFA for this alloy seems to be overestimated by Wang et al. [27],in which the Dc was claimed to be 2 mm. In addition, our findingdoes not support their explanation for the BMG formation with ametastable eutectic concept.

It has been well recognized that the short-range order (SRO) inthe undercooled liquid is a structurally predominant factor for theGFA of alloys. Strong interatomic interactions between unlike-atomic pairs and ordering of atomic arrangement allow the under-cooled melt to attain a high resistance against the solute partitionwhich is necessary for the formation of crystal nuclei. Meanwhile,the structure of metallic glasses also directly correlates with thestructure of the eutectic liquids [28,29]. As shown by Sadoc et al.[30], regarding the structure of CuxZr100�x (x = 33,46,60) glassesstudied by using EXAFS, the distribution of Cu–Cu pairs remains al-most constant in the whole range of concentration and has somesimilarity to the SRO in crystalline Cu10Zr7, while the distributionof the unlike-atom pairs varies drastically, becoming more asym-metric with increasing Cu content. In the meantime, the chemicalordering increases. Therefore, it was suggested that the Cu atomspossess a well-defined local environment and as such would forma fixed matrix in which the Zr atoms are constrained to fit.

In this sense, one can speculate that for the Cu–Zr–Ti ternary al-loys with less than about 10 at.% Ti, which are only related to the e3

reaction, the atomic configuration in the alloy melt resemble thecase of Cu–Zr binary alloys covered by the e3 reaction. It is ex-pected that the Cu10Zr7-like clusters are probably present in theundercooling liquid. Thus, upon cooling the melt the glass forma-tion competes with the crystallization to form the Cu10Zr7 phase.In other words, the GFA depends on whether or not the nucleationof Cu10Zr7 phase can be suppressed. With respect to the Cu–Zrpairs, the interaction between Cu and Ti is somewhat weak, asindicated by the more negative heat of mixing for Cu–Zr

Table 2Eutectic reaction and BMG formation in the Cu–Zr binary system

No. Eutectic reaction BMG-formingcomposition (at.%)

Dc (mm) Refs.

e1 L ? (Cu) + Cu9Zr2 Nonee2 L ? Cu8Zr3 + Cu10Zr7 Cu64.5Zr35.5 2 [19]

Cu64Zr36 2 [24]Cu60Zr40 1.5 [25]

e3 L ? Cu10Zr7 + CuZr Cu54Zr46, Cu52Zr48 1 [8]Cu51Zr49 �1.5 This work

e4 L ? CuZr + CuZr2 Cu46Zr54 2 [7]e5 L ? CuZr2 + (bZr) None

(DHmix = �23 kJ/mol) than that for Cu–Ti (DHmix = �9 kJ/mol)[31]. Additionally, the heat of mixing for Zr–Ti is zero [31]. As a re-sult, a small amount of Ti addition into the Cu–Zr reduces the de-gree of order in the Cu10Zr7-like clusters in the undercooled liquid.Furthermore, considering the topological effect, the atomic radiusof Ti (rTi = 0.146 nm) is intermediate between Cu (rcu = 0.128 nm)and Zr (rZr = 0.160 nm). As substitutional solutes for Zr, the Ti addi-tion at an appropriate concentration would help improve theatomic pack efficiency of the liquid and restrain atomic mobility.As proposed in Ref. [32,33], the GFA is favored by having a moreuniform separation in the atomic sizes, as well as a wider atomicsize distribution range. Therefore, in contrast to the correspondingTi-free binary alloys, the liquid of Ti-containing alloys is more sta-ble, manifested by a deeper eutectic reaction. As demonstrated byour experiments, the Ti addition in the Cu–Zr alloys do play a sig-nificant role of depressing the liquidus and solidus temperatures,lowering the solidus temperature by as much as about 27–53 Kwith respect to the Cu10Zr7–CuZr binary subsystem.

Furthermore, it should be mentioned that the composition re-gion we discovered to form the BMGs with Dc = 4 mm is com-pletely different from the Cu60Zr30Ti10 previously reported byInoue0s group. It indicates that more than one BMG-forming com-position zones are present in the Cu–Zr–Ti ternary system, relatedto different invariant or univariant deep-eutectic reactions.

5. Conclusions

1. Our experiments demonstrated that in the Cu–Zr–Ti ternarysystem, the Cu52Zr40Ti8 composition is not located at a ternaryinvariant eutectic point (L ? Cu10Zr7 + CuZr + Cu2ZrTi), as previ-ously predicted using the CALPHAD calculations. Solidificationof the Cu52Zr40Ti8 liquid upon cooling mainly undergoes a(L ? Cu10Zr7 + CuZr) univariant eutectic reaction. With respectto the Cu–Zr binary alloys correlated with this reaction, the Tiaddition results in a significant depression of the melting onsettemperature and the liquidus of the alloys, about 55 and 27 K,respectively.

2. Alloys around the Cu52Zr40Ti8 exhibit a good glass-forming abil-ity. Monolithic glassy rods of 4 mm in diameter can be fabri-cated under copper mold casting at five compositions,Cu52Zr42Ti6, Cu51Zr42Ti7, Cu51Zr41Ti8, Cu52Zr40Ti8 andCu51Zr40Ti9. Compared with the corresponding Cu–Zr binaryalloys, the presence of Ti significantly promotes the glass for-mation. For the Cu–Zr binary alloys correlated with the eutecticreaction of (L ? Cu10Zr7 + CuZr), the optimal composition forglass formation is off-eutectic, at Cu51Zr49 with a capability toform a near 1.5-mm-diameter glassy rod. The compositionregion we currently discovered in the Cu–Zr–Ti ternary systemto form the BMGs with Dc = 4 mm is independent of that previ-ously reported by Inoue0s group, at Cu60Zr30Ti10, indicating thatmore than one BMG-forming composition zones are present inthis system. Each of these zones is directly correlated with theirrespective univariant eutectic reaction [34].

Acknowledgment

This work was supported by the National Natural Science Foun-dation of China under Grants Nos. 50671105 and 50621091.

References

[1] X.H. Lin, W.L. Johnson, J. Appl. Phys. 78 (1995) 6514.[2] A. Inoue, W. Zhang, T. Zhang, K. Kurosaka, Acta Mater. 49 (2001) 2645.[3] J.C. Lee, Y.C. Kim, J.P. Ahn, S. Lee, B.J. Lee, Appl. Phys. Lett. 84 (2004)

2781.[4] J. Das, M.B. Tang, K.B. Kim, R. Theissmann, F. Baier, W.H. Wang, J. Eckert, Phys.

Rev. Lett. 94 (2005) 205501.

C.-L. Dai et al. / Journal of Non-Crystalline Solids 354 (2008) 3659–3665 3665

[5] P. Wesseling, T.G. Nieh, W.H. Wang, J.J. Lewandowski, Scripta Mater. 51 (2004)151.

[6] A. Inoue, W. Zhang, T. Tsurui, A.R. Yavari, A.L. Greer, Philos. Mag. Lett. 85 (2005)221.

[7] D.H. Xu, G. Duan, W.L. Johnson, Phys. Rev. Lett. 92 (2004) 245504.[8] S.W. Lee, M.Y. Huh, E. Fleury, J.C. Lee, Acta Mater. 54 (2006) 349.[9] C.L. Dai, H. Guo, Y. Shen, Y. Li, E. Ma, J. Xu, Scripta Mater. 54 (2006) 1403.

[10] Y. Shen, E. Ma, J. Xu, J. Mater. Sci. Technol. 24 (2008) 149.[11] P. Jia, H. Guo, Y. Li, J. Xu, E. Ma, Scripta Mater. 54 (2006) 2165.[12] C.G. Woychik, T.B. Massalski, Z. Metallkd. 79 (1988) 149.[13] R. Arroyave, T.W. Eagar, L. Kaufman, J. Alloy Compd. 351 (2003) 158.[14] D. Wang, H. Tan, Y. Li, Acta Mater. 53 (2005) 2969.[15] D. Arias, J.P. Abriata, in: T.B. Massalski, H. Okamoto, P.R. Subramanian (Eds.),

Binary Alloy Phase Diagrams, ASM International, 1990, p. 1511.[16] D.V. Louguine-Lugin, A.D. Setyawan, H. Kato, A. Inoue, Appl. Phys. Lett. 88

(2006) 251902.[17] D. Schryvers, G.S. Firstov, J.W. Seo, J. VanHumbeeck, Y.N. Koval, Scripta Mater.

36 (1997) 1119.[18] H. Tan, Y. Zhang, D. Ma, Y.P. Feng, Y. Li, Acta Mater. 51 (2003) 4551.[19] D. Wang, Y. Li, B.B. Sun, M.L. Sui, K. Lu, E. Ma, Appl. Phys. Lett. 84 (2004)

4029.

[20] H. Ma, Q. Zheng, J. Xu, Y. Li, E. Ma, J. Mater. Res. 20 (2005) 2252.[21] W.J. Boettinger, in: B.C. Kear, B.C. Giessen, M. Cohen (Eds.), Rapidly Solidified

Amorphous and Crystalline Alloys, Elsevier Science Pub. Co., 1982, p. 15.[22] J.D. Wilde, L. Froyen, S. Rex, Scripta Mater. 51 (2004) 533.[23] J.D. Wilde, L. Froyen, V.T. Witusiewicz, U. Hecht, J. Appl. Phys. 97 (2005)

113515.[24] D.H. Xu, B. Lohwongwatana, G. Duan, W.L. Johnson, C. Garland, Acta Mater. 52

(2004) 2621.[25] A. Inoue, W. Zhang, Mater. Trans. 45 (2004) 584.[26] K. Hajlaoui, A.R. Yavari, B. Doisneau, A. LeMoulec, W.J. Botta, F.G. Vaughan, A.L.

Greer, A. Inoue, W. Zhang, A. Kvick, Scripta Mater. 54 (2006) 1829.[27] W.H. Wang, J.J. Lewandowski, A.L. Greer, J. Mater. Res. 20 (2005) 2307.[28] A.R. Yavari, Nat. Mater. 4 (2005) 2.[29] L.L. Shi, J. Xu, E. Ma, Acta Mater. 56 (2008) (in press).[30] A. Sadoc, Y. Calvayrac, A. Quivy, M. Harmelin, A.M. Flank, J. Non-Cryst. Solids

65 (1984) 109.[31] F.R. de Boer, R. Boom, W.C.M. Mattens, A.R. Miedema, A.K. Niessen, Cohesion in

Metals, Elsevier B.V., 1988. p.9.[32] D.B. Miracle, W.S. Sanders, O.N. Senkov, Philos. Mag. 83 (2003) 2409.[33] F.Q. Guo, S.J. Poon, G.J. Shiflet, J. Appl. Phys. 97 (2005) 013512.[34] C.L. Dai, J.W. Deng, Z.X. Zhang, J. Xu, J. Mater. Res. 23 (2008) 1249.