Enhancing bulk metallic glass formation in Ni–Nb–Sn-based ...NbSn. In the map, the open, half open, and solid symbols represent crystalline, partially amorphous, and fully amorphous

Enhancing bulk metallic glass formation in Ni–Nb–Sn-basedalloys via substitutional alloying with Co and Hf

Li Zhang, Mu-Jin Zhuo, and Jian Xua)

Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academyof Sciences, Shenyang 110016, China

(Received 11 June 2007; accepted 12 September 2007)

Bulk metallic glasses have been formed over a fairly wide composition range(54–62 at.% Ni, 32–36 at.% Nb, and 3–11 at.% Sn) in the Ni–Nb–Sn ternary system.Partial substitution of Co for Ni and Hf for Nb improves the glass-forming ability,eventually leading to 4 mm glassy rods at the Ni56Co3Nb28Hf8Sn5 composition. Thepositive effects of these alloying elements have been explained based on a systematicmonitoring of the amount and morphology of the competing crystalline phases as afunction of the Co and Hf contents.

I. INTRODUCTION

The Ni–Nb binary system is well known to exhibitgood glass-forming ability (GFA).1,2 Bulk metallicglasses (BMGs) with a critical size (Dc) of about 1 mmdiameter can be fabricated using conventional coppermold casting in this system.3,4 Starting from the binaryBMG, the GFA was improved through a simple multi-component substitution approach to increase the Dc to 2to 3 mm, such as in Ni(Co, Cu)–Nb(Ti, Zr, Hf)5–7 and Ni–Nb(Ta)8 systems. Another route is to add Sn into thebinary base alloy. The Ni–Nb–Sn ternary system con-tains a eutectic (L → Ni3Nb + Ni6Nb7 + Ni2NbSn) muchdeeper than that in the Ni–Nb binary [the eutectic tem-perature (Teut) for the ternary is about 85 K lower thanthat for the binary].9 The GFA is thus increased, as in-dicated by the successful fabrication of fully glassy rect-angular strips up to 3 mm in casting thickness.9 Theconstituents exhibit strong attractive interactions10: re-cently the density of the binary and the ternary alloy wasmeasured to be about 6% and 9% higher, respectively,than that for the ideal mixing of elements.

Among the BMGs, Ni-based alloys are particularlyinteresting because of their ultrahigh fracture strength(even higher than 3 GPa), high thermal stability, goodcorrosion resistance, and excellent wear resistance.5,11–14

Unfortunately, its Dc remains15–17 significantly belowthose of the other BMGs based on other engineeringmetals such as Fe,18–20 Cu,21–23 Mg,24 let alone thosebased on noble metals (Pd25 and Pt26) or early transitionmetals (Zr27,28 or rare earth29,30). It is thus important to

further explore the effects of alloying elements onthe GFA of Ni-based BMGs. Since the GFA of the alloyin a given system can be strongly composition-dependent,31,32 it has been a challenging problem to op-timize the easy glass-forming composition in a systemwith more than three constituents. Recently, a new ap-proach to pinpoint the best glass former in three dimen-sions (3D) composition space was established.24,33

Adopting this “3D pinpointing approach,” record-sizeBMGs in Mg-based24,34 and Cu-based22 quaternary sys-tems have been discovered. Such a methodology is yet tobe applied to Ni-based systems.

The purpose of this paper is threefold. First, the com-position-dependent GFA in the Ni–Nb–Sn ternary sys-tem is revisited, starting from Choi-Yim’s alloy(Ni59.5Nb33.6Sn6.9).9 We will demonstrate a wide rangeof compositions within which 2-mm BMGs can be pro-duced. Second, Co was selected to substitute for Ni, andthe best BMG-forming alloy was located in the Ni–Co–Nb–Sn quaternary system using the “3D pinpointing ap-proach.” Hf was then used to partially substitute for Nbin the quaternary alloy to further improve GFA. The thirdaspect, which is one of the highlights of this work, is asystematic characterization of the microstructure (phaseidentification, composition, morphology, and volumefraction) of the competing crystalline phases as a func-tion of the concentration of the substituting elements.The detailed information regarding phase selection andmorphology is used to establish the correlation betweenthe BMG-forming composition and the eutectic reactionthat the liquid undergoes during cooling with and withoutthe substituting elements, and in particular to uncover theexact role that the Co and Hf substitutions played inpromoting the GFA. This comprehensive monitoring ofphases/microstructure sets this work apart from most pre-vious studies, where the effects of substituting elements

a)Address all correspondence to this author.e-mail: [email protected]

DOI: 10.1557/JMR.2008.0084

J. Mater. Res., Vol. 23, No. 3, Mar 2008 © 2008 Materials Research Society688

were often explored on trial-and-error basis or reportedwithout explanation.5–8

II. EXPERIMENTAL

Elemental pieces with purity better than 99.9 wt%were used as starting materials. The master alloys withthe nominal composition (in atomic percentage) wereprepared by arc melting under a Ti-gettered argon atmos-phere in a water-cooled copper hearth. The alloy ingotsof 25 g in weight were melted several times to ensurecompositional homogeneity. For the bulk samples with adiameter between 2 and 4 mm, the master alloy wasremelted in a quartz tube using induction melting andinjected in a purified inert atmosphere into the coppermold that has internal rod-shaped cavities of about50 mm in length.

The cross-sectional surfaces of the as-cast rods wereanalyzed by x-ray diffraction (XRD) using a D/max 2400diffractometer (Rigaku, Tokyo, Japan) with monochro-mated Cu K� radiation. The morphology of cross sectionsfor the arc melted ingots was observed by using a scan-ning electron microscope (SEM; Philips XL30, TheNetherlands). The local compositions were semiquan-titatively determined using an energy-dispersivex-ray spectrometer (EDX) attached to the SEM. Themicrostructures were also analyzed using JEM-2010(JOEL, Tokyo, Japan) transmission electron microscope(TEM). TEM samples were mechanically polished andthen thinned by ion milling using purified argon.

The glass transition, crystallization, and melting be-havior of as-cast glassy samples were investigated usinga Netzsch differential scanning calorimeter (DSC 404 CPegasus, NETZSCH-Gerätebau GmbH, Bayern, Ger-many) with alumina container at a heating and coolingrate of 20 K/min under flowing purified argon after beingevacuated to a vacuum of ∼10−3 Torr. A preheating runup to the temperature between glass-transition tempera-ture (Tg) and onset crystallization temperature (Tx1) wasused to remove the irreversible relaxation. A second rununder identical conditions was used to determine thebaseline after each run. To confirm the reproducibility ofthe experimental results, at least three samples have beenmeasured for each composition. All the Tg and Tx1 meas-urements were reproducible within the error of ±2 K. Theheat of crystallization �Hx for the glassy phase was de-termined by integrating the area under the DSC curve.

III. RESULTS

A. Compositional dependence of the BMGformation in the Ni–Nb–Sn ternary system

Figure 1 displays the BMG-forming composition mapfor 2-mm-diameter as-cast rods. The selected triangular

region is enclosed by three compounds Ni3Nb, Ni6Nb7,and Ni2NbSn. In the map, the open, half open, and solidsymbols represent crystalline, partially amorphous, andfully amorphous phases, respectively. The BMG-formingcompositions for Dc � 2 mm fall within the region of 54to 62 at.% Ni, 32 to 36 at.% Nb, and 3 to 11 at.% Sn. Forcomparison, Choi-Yim’s alloy is marked with a dia-mond, which is located at the boundary of our zone forDc � 2 mm. Though the 2-mm BMG formation zone iswide, fully amorphous samples were not obtained for the3-mm-diameter rods at any of these 23 compositions(only a small amount of glassy phase appeared at theouter layer of the rods).

Figure 2(a) shows the XRD patterns taken from thecross-sectional surface of as-cast 2-mm BMGs for sev-eral representative compositions, including Ni62Nb35Sn3,Ni59Nb35Sn6, and Ni57Nb34Sn9. No Bragg peaks fromany crystalline phase are detectable, but a broad diffusivediffraction maximum at 2� � 35 to 50°, indicating theformation of monolithic metallic glasses. The corre-sponding DSC scans are given in Fig. 2(b). In all cases,sharp exothermic crystallization peaks are observed, aswell as clear endothermic events associated with theglass transition. For all these glasses, crystallization oc-curs through multiple steps. The glass-transition tem-perature Tg and the onset temperature of the first exo-thermic peaks Tx1 are marked with arrows on each curve.The measured thermal properties are listed in Table I.Note that both Tg and Tx1 shift to a lower temperatureas the Sn content in the alloy is increased, from 890 and924 K at 3 at.% Sn down to 868 and 906 K at 9 at.% Snfor Tg and Tx1, respectively. The width of the super-cooled liquid region (�Tx � Tx1 − Tg) for these BMGs is

FIG. 1. BMGs formation zone with Dc � 2 mm in the Ni–Nb–Snternary system, near the eutectic related to Ni3Nb, Ni6Nb7 andNi2NbSn. Chio-Yim’s alloy, Ni59.5Nb33.6Sn6.9, is marked as a diamondfor comparison.

L. Zhang et al.: Enhancing bulk metallic glass formation in Ni–Nb–Sn-based alloys via substitutional alloying with Co and Hf

J. Mater. Res., Vol. 23, No. 3, Mar 2008 689

insensitive to the composition variation, remainingaround 35 K. Compared with the Ni(Pd)–P (Tx1 ≈680 K),17 Ni–B(Tx1 ≈ 810 K),35 Ni–(Zr,Ti)–(Al, Si, Cu, Pd) (Tx1 � 800 to 850 K)16,36,37 based BMGs,Ni–Nb–Sn-based BMGs exhibit higher thermal stability.

Figure 3(a) displays the DSC curves during heatingand cooling in the temperature range near the meltingtemperature for the Ni65−xNb35Snx (3 � x � 11) seriesalloys (the Nb content is fixed at 35 at.%). This series ofcompositions is marked as a dash line in Fig. 1. As seenfrom Fig. 3(a), the onset temperature of melting (Tm,marked with an arrow on each curve) for all the alloys isnearly the same, indicating that alloy melts undergo anidentical eutectic reaction during solidification. How-ever, the end temperature of the melting event (defined asthe liquidus temperature, TL, marked with an arrow on each

curve as well) strongly depends on the Sn content, de-creasing dramatically with increasing the Sn concentra-tion. Note that only a single melting event (and a singlesolidifying process) happens at the Ni54Nb35Sn11 com-position, over a rather narrow temperature span of about19 K. This observation suggests that this composition canbe approximately defined as the ternary invariant eutecticpoint for the reaction (L → Ni3Nb + Ni6Nb7 + Ni2NbSn)(see below). We have to emphasize that the eutectic pointwas believed to be at Ni60Nb29Sn11 in Ref. 9. The DSCcurve for this alloy is revisited, as shown as Fig. 3(b).Three main events are clearly visible during either melt-ing or solidifying, indicating that this is not a eutecticcomposition. It is explained later in this article that thesolidification pathway of this alloy actually follows thesequence: primary crystallization of Ni3Nb or Ni2NbSn(so far uncertain), then univariant eutectic (Ni3Nb +Ni2NbSn), and finally ternary invariant eutectic (Ni3Nb +Ni6Nb7 + Ni2NbSn) reaction.

DSC measurements for the melting behavior were car-ried out for most of the alloys within the investigatedcomposition range shown in Fig. 1. Using the data of Tm

and TL from 25 compositions, 3D liquidus and solidussurfaces of the system are created, as drawn in Fig. 4.The Tm of all the alloys is nearly the same within ex-perimental error, suggesting that all of these alloys un-dergo an identical ternary eutectic reaction. The invarianteutectic temperature Teut was determined to be 1366 ± 2 K.It is in agreement with the value (Teut � 1363 K) fromRef. 9. In contrast, the liquidus surface exhibits a fairlysteep descent toward the eutectic point, reflecting a rela-tively deep eutectic feature of the system. The slope ofthe liquidus, starting from the melting point of Ni3Nb at1672 K, is 17.2 K/at.%, steeper than the 14.2 K/at.% forthe binary eutectic (L → Ni3Nb + Ni6Nb7). The optimalBMG-forming composition zone is drawn as a red circleon the bottom projection in Fig. 4. It is seen that this zoneis off-eutectic toward the Sn-lean side, the side with asteeper liquidus slope. Such off-eutectic glass-formingregions have been explained previously in terms of thecoupled eutectic zone concept.38,39

The reduced glass-transition temperature Trg (Trg �Tg/TL) is calculated for the typical ternary BMGs andlisted in Table I as well, together with the TL for eachalloy. Although the Trg increased slightly from 0.58 at3 at.% Sn up to 0.60 at 9 at.% Sn, there is no significantdifference for the GFA among these alloys.

B. Optimizing the GFA in the Ni–Co–Nb–Snquaternary system using the “3Dpinpointing approach”

Starting from the BMG-forming composition zone inthe ternary, Co was introduced to substitute for Ni toenhance the GFA. In light of the similarity in chemistryand atomic size between Ni and Co, the Ni–Co–Nb–Sn

FIG. 2. (a) XRD patterns taken from the cross-sectional surface. (b)DSC scans of several representative as-cast BMGs in the 2-mm glassforming zone of the ternary alloy system.


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quaternary was treated as the (Ni, Co)–Nb–Sn pseudo-ternary. The optimal BMG-forming composition was lo-cated in the 3D composition space following the stepwisenavigation protocol using the “3D pinpointing ap-proach.” Several consecutive compositional planes wereexamined, each with a fixed Co-to-Ni ratio expressed asNi1−xCox (x � 0.05, 0.1, 0.125, 0.15, respectively). Fig-ures 5(a)–5(d) display the BMG-forming compositionmap for 3-mm-diameter as-cast rods. The GFA was im-proved even with minor Co substitution, such that at x �0.05 the Dc of two Co-containing alloys reached 3 mm.This size remains unchanged when the Co-to-Ni ratioincreased to x � 0.15, but to maintain this Dc it wasnecessary to simultaneously tune the concentrations ofNb and Sn, moving away from the Nb-rich side when xwas increased. With the composition interval of 1 at.%, atotal of eight alloy compositions were found to havenearly equivalent GFA (Dc � 3 mm). The optimal ratioof Co to Ni is located at x � 0.125, see Fig. 5(c).

Figures 6(a) and 6(b) display the XRD patterns andDSC scans, respectively, for the 3-mm-diameter as-castBMGs selected from each plane. Similar to the Co-freeternary BMGs, the crystallization was completed throughmultiple steps in all cases, showing at least three sharp

TABLE I. Thermal properties of Ni–Nb–Sn based BMGs fabricated using copper mold casting, determined using DSC at a heating rate of20 K/min.

Composition Dc (mm) Tg (K) Tx1 (K) �Tx (K) �Hx (kJ/mol) Tm (K) TL (K) Trg

Ni62Nb35Sn3 2 890 924 34 7.1 1366 1522 0.58Ni59Nb35Sn6 2 890 924 34 8.7 1366 1497 0.59Ni57Nb34Sn9 2 868 906 38 7.0 1366 1439 0.60Ni56Co3Nb36Sn5 3 872 916 44 7.1 1368 1462 0.60Ni54Co6Nb34Sn6 3 872 917 41 8.7 1370 1465 0.60Ni52.5Co7.5Nb34Sn6 3 872 920 48 7.2 1369 1466 0.60Ni51Co9Nb33Sn7 3 872 922 50 8.5 1372 1467 0.60Ni56Co3Nb28Hf8Sn5 4 864 919 55 6.4 1385 1415 0.61

FIG. 3. (a) DSC scans of the alloy series Ni65−xNb35Snx (3 � y � 11).(b) Ni60Nb29Sn11 near their melting temperatures.

FIG. 4. Liquidus and solidus surfaces for the Ni–Nb–Sn system,measured from DSC curves. The eutectic composition and the com-position range for the 2-mm BMGs (red zone) are marked.



exothermic peaks in the DSC scans, as seen in Fig. 6(b).With the increasing of Co content, the second crystalli-zation peak around 950 K fades away and finally disap-pears when Co content reached 9 at.%. It was noticed thatthe supercooled liquid region, �Tx, for these quaternaryBMGs is ∼10 K larger than those of the ternary ones.This is consistent with the finding that the substitutionalalloying with Co results in an enhancement of the GFA.

Figure 7 shows the DSC curves during heating andcooling near and above their melting temperature for theCo-containing BMGs, corresponding to the alloys in Fig.6. A curve for a ternary alloy Ni60Nb33Sn7 (x � 0) isplotted together to compare with the corresponding qua-ternary alloy Ni51Co9Nb33Sn7. For all of the quaternaryalloys, the melting or solidifying curves contain threeevents, indicating that none of them is situated at eitheran invariant eutectic point or in a univariant eutecticgrove. Compared with the ternary, the Tm of the quater-nary alloys remains unchanged at around 1370 K. The TL

of the alloys with 3 to 9 at.% Co is around 1465 K.However, with respect to the Ni60Nb33Sn7 ternary, the TL

of the Ni51Co9Nb33Sn7 quaternary is lowered by 24 K,from 1491 down to 1467 K, providing the evidence thatalloying with Co stabilizes the liquid. The thermal prop-erties obtained form DSC measurements, including Tg,

Tx1, �Tx, �Hx, Tm, TL, and calculated Trg values, are alsolisted in Table I for comparison. The Trg of Co-containing quaternary alloys are nearly the same at 0.60.

C. Effect of Hf substitution for Nb in theNi–Co–Nb–Sn quaternary on the GFA

To further improve the GFA of the Ni–Co–Nb–Snquaternary alloys, Hf was introduced to partially substi-tute for Nb in the Ni56Co3Nb36Sn5 alloy. This leads to aseries of alloys, Ni56Co3Nb36−yHfySn5 (0 � y � 10),which were cast into 4 mm diameter rods. Figure 8 dis-plays the XRD patterns taken from the cross-sectionalsurface of these as-cast alloys. As the Hf content wasincreased in the range of 2 � y � 6, the intensities of thediffraction peaks from the Ni3Nb and Ni6Nb7 crystallinephases were gradually reduced, indicating that the for-mation of these intermetallics was frustrated due to thepresence of Hf. When the content of Hf was increased to8 at.%, diffraction peaks from the crystalline compoundis no longer detectable beside the broad diffusive diffrac-tion maximum at 2� � 35 to 50°, indicating that a mono-lithic metallic glass can be formed at this size. How-ever, on further increasing the Hf content up to 10 at.%,the GFA was degraded again. New diffraction peaksfrom unidentified crystalline phase (not consistent with

FIG. 5. BMG formation in the (Ni1−xCox)–Nb–Sn system. (a–d) Show the composition ranges for BMG formation (Dc � 3 mm) for several ratiosof Ni1−xCox. (a) x � 0.05, (b) x � 0.1, (c) x � 0.125, and (d) x � 0.15.



Ni3Nb, Ni6Nb7, or Ni2NbSn) appeared. This indicatesthat at sufficiently high Hf content new crystalline phasesemerge to compete with the glass formation.

Figure 9 shows a DSC scan of the Ni56Co3Nb28Hf8Sn5

as-cast 4 mm diameter BMG, together with a scan for the1 mm diameter BMG of the identical composition forcomparison. Unlike the Ni–Co–Nb–Sn quaternaryglasses shown in Fig. 6(b), there are two sharp crystal-lization peaks in the scan. The width of the supercooledliquid region, �Tx, is extended to 55 K, wider by 10 Kwith respect to the Hf-free Ni–Co–Nb–Sn BMGs. The�Hx value of 1- and 4-mm BMGs is determined to be 6.5and 6.4 kJ/mol, respectively. This agreement within ex-perimental error confirms complete glass formation inthe 4-mm-diameter as-cast rod. The inset in Fig. 9 showsDSC curves associated with the melting and solidifyingevents for the Ni56Co3Nb28Hf8Sn5 BMG. In both cases,

two events are observed, indicating that the alloy is at anoff-eutectic composition. The Tg, Tx1, �Tx, Tm, TL, Trg,and �Hx obtained from DSC measurements are includedin Table I as well. In contrast to the Ni–Co–Nb–Sn qua-ternary glasses, the Tm of the Hf-alloyed BMG is higherby 15 K, but TL is lowered by 40 K, resulting in a slight

FIG. 6. (a) XRD patterns and (b) DSC scans of representative as-castBMGs for the four ratios of Ni1−xCox (x � 0.05, 0.1, 0.125, 0.15)corresponding to the four planes shown in Fig. 5.

FIG. 7. High-temperature DSC scans of the representative BMGformers corresponding to Fig. 6, together with the Co-free alloyNi60Nb33Sn7 for comparison.

FIG. 8. XRD patterns of as-cast rods 4 mm in diameter for the alloyseries Ni56Co3Nb36−yHfySn5 (y � 0, 2, 4, 6, 8, 10).



increase of Trg value to 0.61. The increase of �Tx and Trg

is consistent with an enhancement of the GFA of thealloy because of Hf substitution.

D. BMG formation correlated with eutecticreaction and the effects of alloying elements onsolidification behavior

The cooling rate for the arc melted BMG-forming al-loy ingot was estimated to be around 1 K/s.40 When thiscooling rate is used, the solidification microstructure ofthe ingot can be directly correlated with the phase selec-tion when the alloy melt crystallizes before freezing intoa glass, i.e., its nonequilibrium solidification pathwayprior to glass formation. In the current work, the phasespresent and their morphology in the arc melted Ni–Nb–Sn alloys with or without Co and Hf are characterized indetail to reveal the role Co and Hf played in affecting thesolidification pathway.

Figure 10 illustrates the XRD patterns taken from thecross-sectional surfaces of the arc melted ingots, forNi59Nb35Sn6, Ni51Co9Nb33Sn7, and Ni56Co3Nb28Hf8Sn5

BMG-forming alloys. For the ternary alloy, three inter-metalllic phases are identified: orthorhombic-Ni3Nb (�-Cu3Ti-type, a � 0.5114 nm, b � 0.4244 nm, c � 0.453nm), hexagonal-Ni6Nb7 (W6Fe7-type, � phase,a � 0.4893 nm, c � 2.664 nm), and face-centered cubic(fcc) Ni2NbSn (BiF3-type, a � 0.6160 nm). The Co-containing quaternary alloy ingot shows the same phases,but it was noticed that the intensity of diffraction peaksfrom Ni3Nb in Ni51Co9Nb33Sn7 is reduced relative tothat of the Ni59Nb35Sn6 ternary. This suggests that Co

has the effect of retarding the formation of Ni3Nb. How-ever, the substitution of Ni with Co made no difference tothe phases participating in the ternary eutectic reactionduring cooling (see below). In contrast, an unidentifiedphase appears in the pattern of the Hf-containing quinaryalloy, besides the Ni6Nb7 and Ni2NbSn phases.

Figures 11(a), 11(c), and 11(e) show the backscatteredSEM (BSEM) images taken from the cross-sectional sur-face of arc melted Ni59Nb35Sn6, Ni51Co9Nb33Sn7, andNi56Co3Nb28Hf8Sn5 alloys. Figures 11(b), 11(d), and11(f) are the high-magnification images for the boxedareas in Figs. 11(a), 11(c), and 11(e), respectively. Thesesurfaces were polished for SEM observations, which re-vealed that the microstructure was uniform throughoutthe sample. The chemical compositions of the phasesexhibiting different contrast were determined using EDXanalysis from the average of at least five measurements,as listed in Table II. The results are consistent with thephase identifications from the XRD experiments. Asmarked (A1) in Fig. 11(a), the acicular (darker) phase(about 5–10% in volume fraction) distributed in the ma-trix, with 10–50 �m in width and 0.1–1 mm in length, isthe Ni3Nb phase precipitated as the primary phase. Therest (the matrix) was found to contain three phases whenobserved at high magnification, marked as A1, B1, andC1 in Fig. 11(b). They are, respectively, the acicularNi3Nb (finer than primary phase), Ni6Nb7 (gray), anddendritic Ni2NbSn (lighter). Such a microstructure istypical of the morphology observed for nonfacetted–facetted eutectics (such as Al–Si or Fe–C binary).41 Inaddition, it was noticed that Sn has a large solubility(∼7%) in the Ni6Nb7 binary compound.

The arc melted Ni51Co9Nb33Sn7 quaternary alloy ex-hibits a similar microstructure, as shown in Figs. 11(c)

FIG. 9. DSC scans of the best glass former Ni56Co3Nb28Hf8Sn5 at adiameter of 1 and 4 mm. The inset shows the melting behavior of thisalloy.

FIG. 10. XRD patterns of the arc melted alloys Ni59Nb35Sn6,Ni51Co9Nb33Sn7, and Ni56Co3Nb28Hf8Sn5.



and 11(d). The acicular Ni3Nb phase remains as the pri-mary phase. However, the eutectic matrix was signifi-cantly refined when compared with the ternary case.Moreover, it was found that the Co exhibits a larger

solubility in all the three phases, Ni3Nb, Ni6Nb7, andNi2NbSn (see Table II), marked as A2, B2, and C2 in Fig.11(d), respectively. Apparently, the Co dissolved in thethree phases retards the coupled growth of the eutectic

FIG. 11. Backscattered SEM images of the arc melted alloys. (a) Ni59Nb35Sn6, (c) Ni51Co9Nb33Sn7, and (e) Ni56Co3Nb28Hf8Sn5. (b), (d), and(f) are the high-magnification view of (a), (c), and (e), respectively.



colony without altering the ternary invariant eutectic re-action (L → Ni3Nb + Ni6Nb7 + Ni2NbSn).

Different from the ternary and quaternary, the primaryphase Ni3Nb is absent in the Hf-containing quinary alloy(Ni56Co3Nb28Hf8Sn5), as demonstrated in Figs. 11(e)and 11(f). The eutectic microstructure was so refined thatit is unable to resolve the layered structure in the SEMimages. Consequently, the Hf addition not only sup-presses the formation of the Ni3Nb phase, but also prob-ably plays a role in slowing down the coupled growth ofthe eutectic structure.

Figures 12(a) and 12(b) display the bright-field TEMimages and selected area electron-diffraction (SAED)patterns of the eutectic matrix in the arc meltedNi59Nb35Sn6 and Ni56Co3Nb28Hf8Sn5, respectively. Forthe Ni59Nb35Sn6 ternary, the three crystalline phases inthe eutectic matrix are in areas A1, B1, and C1, respec-tively, in Fig. 12(a). They have been identified from theSAED patterns, shown atop Fig. 12(a) from left to right.The SAED pattern for area A1, taken along the [142]axis , is indexed as the orthorhombic Ni3Nbphase with a � 0.512 nm, b � 0.424 nm, and c �0.453 nm. The area B1 is identified to be hexagonalNi6Nb7, based on the pattern taken along the [241] axis.The lattice parameters for this phase were determined tobe a � 0.509 nm and c � 2.799 nm, slightly larger thanthose of standard binary compound (a � 0.489 nm andc � 2.66 nm). This lattice expansion is caused by therelatively large Sn solutes dissolved in this phase. The C1

area is indexed to be fcc Ni2NbSn ternary compoundwith a � 0.616 nm, from its SAED pattern taken alongthe [100] axis. These findings are consistent with theresults from the XRD analysis and SEM observations. Inthe arc melted ingot, the liquid of the BMG-formingalloy undergoes the ternary invariant eutectic reaction(L → Ni3Nb + Ni6Nb7 + Ni2NbSn) as the cooling ratewas not sufficiently high to form the glass.

As seen Fig. 12(b), three areas with different contrastare observed in the bright-field TEM image for the arcmelted Ni56Co3Nb28Hf8Sn5. These regions are marked asA3, B3, and C3, respectively. The corresponding SAEDpatterns are shown in the insets from left to right. The

TABLE II. Chemical composition determined using EDX for different phases in the arc melted Ni59Nb35Sn6, Ni51Co9Nb33Sn7, andNi56Co3Nb28Hf8Sn5 alloys. (The experimental error of the measurement was estimated to be ±1 at.% using a reference.)

Alloy Area Ni (at.%) Co (at.%) Nb (at.%) Hf (at.%) Sn (at.%) Normalized phase

Ni59Nb35Sn6 A1 74 �� 26 �� 0 Ni3NbB1 49 �� 44 �� 7 Ni6Nb7

C1 53 �� 27 �� 20 Ni2NbSnNi51Co9Nb33Sn7 A2 65 9 26 �� 0 (Ni,Co)3Nb

B2 44 11 38 �� 8 (Ni,Co)6Nb7

C2 48 7 25 �� 20 (Ni,Co)2NbSnNi56Co3Nb28Hf8Sn5 A3 69 3 20 8 0 (Ni,Co)3(Nb,Hf)

B3 47 4 42 3 4 (Ni,Co)6(Nb,Hf)7

C3 53 2 12 13 21 (Ni,Co)2(Nb,Hf)Sn

FIG. 12. Bright-field TEM images of arc melted (a) Ni59Nb35Sn6 and(b) Ni56Co3Nb28Hf8Sn5. Insets are the SAED patterns of each markedphase, corresponding to the areas of A1, B1, and C1 in (a), and A3, B3,and C3 in (b) from left to right.



composition of each of these areas measured by EDX isalso summarized in Table II. In contrast to the ternaryalloy, the feature size in the eutectic structure of theHf-containing quinary alloy was significantly refineddown to 50 to 100 nm. For the area A3, the compositioncan be reduced as the (Ni, Co)3(Nb, Hf) phase. However,its structure could not be identified as either the ortho-rhombic Ni3Nb or the hexagonal Ni3Hf (�-(Pd, Rh)3Ta-type, a � 0.5282 nm and c � 2.1392 nm). Thus, thisphase is probably a metastable phase with a compositionquite similar to stoichiometric Ni3Nb but with an un-known structure. It provides strong evidence to revealthat the substitutional alloying with Hf for Nb can effec-tively suppress the formation of the Ni3Nb phase in thealloy. Based on the SAED patterns (taken along the [120]and [100] axial, respectively), the B3 and C3 areas areindexed as the hexagonal Ni6Nb7 and fcc Ni2NbSn, re-spectively. In both cases, Ni and Nb atoms in the phasesare partially substituted by the Co and Hf atoms, respec-tively, as seen in Table II. These findings confirm that theeutectic reaction involved in the Co- and Hf-containingquinary alloy remains [L → Ni3Nb (metastable) +Ni6Nb7 + Ni2NbSn], the same as that in the Ni–Nb–Snternary. The difference is that the coupled growth of thethree phases is markedly limited in the quinary system.This is directly correlated with the GFA increasing in thequinary alloy with respect to the Ni–Nb–Sn ternary.

IV. DISCUSSION

It is well understood that the formation of metallicglass competes with the crystallization of the under-cooled melt. For a given lower-order alloy system, it isnecessary to improve the GFA either by stabilizing theliquid or by destabilizing the competing crystallinephases that are usually intermetallics. To date, it has beena conventional approach to design multicomponent alloysystem via further introducing additional elements. How-ever, the effect of an additional component should becarefully assessed by comparing the best glass former oflower-order alloy system (for example, the ternary alloysystem) with that in the higher-order (e.g., the quater-nary) system. Our previous work for adding Ni to theMg–Cu–Y ternary system33 indicates that before locatingthe optimal BMG-forming composition in the ternarysystem, it is uncertain to judge the effect of the fourthcomponent on the GFA (favorable or adverse). Conse-quently, it remains necessary to locate the best glassformer in a lower-order alloy system through the com-position pinpointing before introducing an additionalcomponent. In the current case, we have uncovered thefull potential of the Ni–Nb–Sn base system. The Dc inthis alloy system is determined to be 2 mm, more than 20alloys form a 2-mm zone located at off-eutectic compo-

sitions. In addition, it was noticed that the compositiondependence of the GFA alloy on the Ni (or Nb) concen-tration is stronger than on the Sn concentration.

Using the center of the 2-mm zone as the startingpoint, Co was introduced to partially substitute for Ni,leading to an increase of Dc from 2 to 3 mm. The positiveeffects of Co may be understood in qualitative terms. TheNi–Co binary system has a nearly zero heat of mixing inthe liquid state without any intermetallics that couldcompete with the glass formation. Meanwhile, Co hasnearly the same atomic radius as Ni (rCo � 0.128 nm andrNi � 0.128 nm). In the resulting Ni–Co–Nb–Sn alloy,Co is supposed to substitute for Ni in the atomic con-figuration of the original Ni–Nb–Sn alloy. For the lattersystem, the primary crystalline phase is orthorhombicNi3Nb. We note that in equilibrium the Co3Nb (hexago-nal MgNi2-type) has a different structure. The addition ofCo thus causes confusion to frustrate the crystallizationprocess and favors glass formation. As seen in Fig. 10,for the Co-containing alloy, the diffraction intensity ofthe competing Ni3Nb phase in the XRD pattern wasfound to decrease.

We observed that GFA increased further after partialsubstitution of Hf for Nb, on the basis of the Ni–Co–Nb–Sn system, leading to the improvement of Dc to 4 mm.The positive effects of Hf can be understood as follows.On the one hand, the Hf addition has stabilized the su-percooled liquid, as evidenced by the reduction of the TL

by as much as ∼40 K. The �Tx increased as well, by10 K with respect to the Hf-free Ni–Co–Nb–Sn BMGs.This is because Hf has a negative heat of mixing in theliquid state with Ni, Co, Sn, and Nb (−42, −35, −35, and−4 kJ/mol,42,43 respectively). The strong interactions fa-vor short-range order and the formation of clusters in theundercooled liquid, lowering its energy.44 On the otherhand, the presence of Hf has clearly frustrated the growthof the competing crystalline phases: the structure of thearc melted master alloy was drastically refined after in-troducing Hf, and the suppression of the competingNi3Nb compound is obvious in the XRD patterns. Thefrustration effect can be attributed to the large atomicradius of Hf (rHf � 0.167 nm), as compared with theothers (rNb � 0.146 nm and rSn � 0.162 nm). Thesolubility of Hf in the competing phases is limited be-cause of the large size difference, as seen in Table II. Forcrystallization upon cooling, Hf atoms have to be redis-tributed. The long-range diffusion of the relatively largeHf atoms is kinetically difficult, slowing down thegrowth rate of the crystal nuclei. As such, Hf has anobvious effect in refining the competing crystallinephases.

In general, the BMG formation can be promoted viasubstitutional alloying with Co and Hf in the Ni–Nb–Snparent ternary alloys. The role that Co and Hf played wasrevealed to correlate with the limitations on eutectic



growth kinetically controlled by diffusion, which is inagreement with the idea set forth earlier by Boettinger.38

V. CONCLUSIONS

A wide composition zone has been identified in whichNi–Nb–Sn ternary alloys can be cast into 2-mm fullyglassy rods. This zone lies at off-eutectic compositionson the side with the steeper liquidus slope. The glass-forming ability can be improved by partial Co substitu-tion for Ni, and the maximum Dc is determined to be3 mm for the pseudo-ternary (Ni, Co)–Nb–Sn alloy sys-tem through the “3D pinpointing approach.” Further sub-stitution of Hf for Nb led to the increase of Dc up to 4 mmat Ni56Co3Nb28Hf8Sn5. Both Co and Hf were found tostabilize the undercooled melt.

A series of XRD, SEM, and TEM experiments havebeen performed to identify the competing phases and tomonitor the effects of the alloying elements on the se-lection and morphology of these competitors to glassformation. This allowed us to conclude that the alloyingsubstitution did not affect the phase selection in the eu-tectic reaction related to glass formation, although thecomponent number was increased in the system, and at-tribute to the beneficial effects of Co and Hf on glass-forming ability to their role in suppressing the formationof the competing intermetallic phase such as Ni3Nb.

ACKNOWLEDGMENTS

The authors gratefully acknowledge the stimulatingdiscussion with Professor E. Ma and Professor X.L. Ma.This research was supported by the National Natural Sci-ence Foundation of China under Grant No. 50621091 andNational Basic Research Program of China (973 Pro-gram) under Contract No. 2007CB613906.

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Enhancing bulk metallic glass formation in Ni–Nb–Sn-based ...NbSn. In the map, the open, half open, and solid symbols represent crystalline, partially amorphous, and fully amorphous