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Materials Chemistry and Physics 125 (2011) 37–45

Contents lists available at ScienceDirect

Materials Chemistry and Physics

journa l homepage: www.e lsev ier .com/ locate /matchemphys

xperimental investigation and thermodynamic calculation of the phasequilibria in the Cu–Mo–Ni ternary system

an Yua, Cuiping Wanga, Yifu Lia, Xingjun Liua,∗, Ryosuke Kainumab, Kiyohito Ishidab

Department of Materials Science and Engineering, College of Materials, and Research Center of Materials Design and Application, Xiamen University, Xiamen 361005, PR ChinaDepartment of Materials Science, Graduate School of Engineering, Tohoku University, Sendai 980-8579, Japan

r t i c l e i n f o

rticle history:eceived 8 March 2010eceived in revised form 12 August 2010ccepted 18 August 2010

eywords:u–Mo–Ni alloys

a b s t r a c t

The phase equilibria at 900 ◦C, 1000 ◦C, 1100 ◦C, 1200 ◦C and 1300 ◦C in the Cu–Mo–Ni system were exper-imentally determined by means of optical microscopy (OM), electron probe microanalyzer (EPMA) andX-ray diffraction (XRD) on the equilibrated alloys. The experimental results firstly found that the fcc-typemiscibility gap exists at 900 ◦C, 1000 ◦C, 1100 ◦C and 1200 ◦C in the Cu–Mo–Ni system, and the solubility ofCu in the MoNi phase at 900 ◦C, 1000 ◦C, 1100 ◦C and 1200 ◦C are about 0.5 at.%, 1.5 at.%, 1.7 at.% and 4.0 at.%,respectively. The as-cast Cu20Mo20Ni60 (at.%), Cu20Mo30Ni50 (at.%), Cu10Mo60Ni30 (at.%), Cu70Mo10Ni60

hase equilibrialectron probeALPHAD

(at.%), Cu20Mo60Ni20 (at.%) and Cu80Mo10Ni10 (at.%) alloys appear the separated macroscopic morpholo-gies, which are caused by the liquid phase separation on cooling, while the as-cast Cu10Mo25Ni65 (at.%),Cu32Mo5Ni63 (at.%) and Cu30.7Mo6.3Ni63 (at.%) alloys show the homogenous microscopic morphologies.On the basis of the experimental data investigated by the present and previous works, the phase equi-libria in the Cu–Mo–Ni system were thermodynamically assessed by using CALPHAD (Calculation ofPhase Diagrams) method, and a consistent set of the thermodynamic parameters leading to reasonable

alcul

agreement between the c

. Introduction

The Cu–Mo base alloys have been widely used as contact mate-ials, protective coatings and ceramic hermetic materials due toheir excellent thermal and electrical conductivities, high corro-ion resistance and superior mechanical properties [1–5]. In theu–Mo binary system, results from previous work [6] indicate thatu and Mo are mutually insoluble at all temperatures in both the

iquid and solid states. The alloying element Ni is usually added tonhance the fusibility and improve the properties of the Cu–Mo–Niernary alloys [7]. Moreover, the properties of alloy materialsre considered to be significantly affected by the microstructurend its composition. In order to well understand the relationshipetween properties and microstructure, accurate information onhe phase equilibria in the Cu–Mo–Ni ternary system is especiallyequired. Although the boundary of fcc single-phase region in theu–Mo–Ni system at 900 ◦C, 950 ◦C, 1000 ◦C and 1100 ◦C deter-ined by Kozlova et al. [8] (the isothermal section at 1000 ◦C is

hown in Fig. 1), the phase equilibria of the Cu–Mo–Ni system at002 ◦C investigated by Kodentsov et al. [9] and a liquidus projec-ion reported by Gupta [10] were available, however, there still areome conflicts, which are summarized as follows: (1) the solubil-

∗ Corresponding author. Tel.: +86 592 2187888; fax: +86 592 2187966.E-mail address: lxj@xmu.edu.cn (X. J. Liu).

254-0584/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.matchemphys.2010.08.068

ated results and experimental data was obtained.© 2010 Elsevier B.V. All rights reserved.

ity of Cu in the MoNi phase measured by Kozlova et al. [8] andKodentsov et al. [9] have a big difference from that in the liquidusprojection reported by Gupta [10] and (2) the critical compositionrange of the Mo-rich liquid miscibility gap in the liquidus projec-tion reported by Gupta [10] is obviously not in accordance with thatin the Cu–Mo binary phase diagram [6]. Therefore, it is importantand necessary to experimentally investigate the phase equilibria inthe Cu–Mo–Ni ternary system.

Obtaining information on the phase equilibria and phasetransformations exclusively from experiments is burdensomeand expensive. The CALPHAD (Calculation of Phase Diagrams)technique, which has been recognized to be an important toolto significantly reduce time and cost during the developmentof materials, can effectively provide a clear guidance for thematerials design [11–13]. In order to develop the thermody-namic database of the Cu base alloys for reliable predictionsof liquidus, phase fraction, equilibrium and non-equilibriumsolidification behavior, etc. in multi-component systems, thepresent authors have made some thermodynamic assessmentsin the Cu base alloy systems [14,15]. As a part of this ther-modynamic database, the thermodynamic description for the

Cu–Mo–Ni ternary system supported by key experiments is of greatessence.

The purpose of the present work is to experimentally investigatethe phase equilibria in the Cu–Mo–Ni ternary system, and to carryout the thermodynamic assessment of the Cu–Mo–Ni ternary sys-

38 Y. Yu et al. / Materials Chemistry an

FC

ts

2

(fn

Ff

ig. 1. The isothermal section at 1000 ◦C with the fcc single-phase region in theu–Mo–Ni system reported by Kozlova [8].

em on the basis of the experimental data and the assessed binaryystems.

. Experimental procedures

Electrolytic copper (99.99 wt.%), molybdenum (99.9 wt.%) and nickel99.95 wt.%) were used as starting materials. Bulk buttons were preparedrom pure elements by arc melting under high purity argon atmosphere using aon-consumable tungsten electrode. The ingots were melted at least six times in

ig. 2. BSE images of the typical ternary Cu–Mo–Ni alloys: (a) Cu30.7Mo6.3Ni63 (at.%) alloyor 45 days, (c) Cu20Mo60Ni20 (at.%) alloy annealed at 1200 ◦C for 20 days, and (d) Cu10Mo

d Physics 125 (2011) 37–45

order to achieve their homogeneity. The sample weight was around 20 g and theweight loss during melting was generally <0.20% of the sample weight. Afterwards,the ingots were cut into small pieces for heat treatments and further observations.

Plate-shaped specimens were wrapped in W foil in order to prevent direct con-tact with the quartz ampoule, and put into quartz ampoule evacuated and backfilledwith argon gas. The specimens were annealed at 1300 ◦C, 1200 ◦C, 1100 ◦C, 1000 ◦Cand 900 ◦C, respectively. The time of the heat treatment varied from several hours toseveral weeks depending on the annealing temperature and the composition of thespecimen. After the heat treatment, the specimens were quenched into ice water.

After standard metallographic preparation, the microstructural observationswere carried out by optical microscopy (OM). The equilibrium compositions of theequilibrated alloys were examined by electron probe microanalyzer (EPMA) (JXA-8100R, JEOL, Japan). Pure elements were used as standards and the measurementswere carried out at 20.0 kV. The X-ray diffraction (XRD) was used to identify thecrystal structure of the constituent phases. The XRD measurement was carried outon a Phillips Panalytical X-pert diffractometer using CuK� radiation at 40 kV and30 mA. The data were collected in the range of 2� from 30◦ to 100◦ at a step widthof 0.0167◦ .

3. Results and discussion

3.1. Phase equilibria and microstructural morphology

BSE (back-scattered electron) images of the typical ternaryCu–Mo–Ni alloys are presented in Fig. 2(a)–(d). Phase identifica-tion was based on the equilibrium composition as measured byEPMA and XRD results. In the Cu30.7Mo6.3Ni63 (at.%) alloy quenchedfrom 1000 ◦C, the separated microstructure (fcc1 + fcc2) (here fcc1and fcc2 phases stand for the fcc Cu-rich and fcc Ni-rich phases)was observed (Fig. 2(a)) and substantiated by the XRD result, as

shown in Fig. 3(a), where only the characteristic peaks of the fccphase were found. Fig. 2(b) shows the three-phase microstruc-ture (MoNi + fcc1 + fcc2) of the Cu20Mo20Ni60 (at.%) alloy annealedat 1000 ◦C for 45 days, and the XRD result is shown in Fig. 3(b),where the characteristic peaks of the MoNi and fcc phases are only

annealed at 1000 ◦C for 45 days, (b) Cu20Mo20Ni60 (at.%) alloy annealed at 1000 ◦C25Ni65 (at.%) alloy annealed at 1300 ◦C for 15 days.

Y. Yu et al. / Materials Chemistry and Physics 125 (2011) 37–45 39

Table 1The equilibrium compositions of the equilibrated alloys determined in the present work.

Temperature (◦C) Alloys (at.%) Annealing time Phase equilibria Composition (at.%)

Phase 1 Phase 2 Phase 3

Phase 1/Phase 2/Phase 3 Mo Ni Mo Ni Mo Ni

900 Cu10Mo25Ni65 60 days MoNi/fcc 51.1 47.9 17.9 73.8Cu20Mo60Ni20 60 days bcc/MoNi/fcc 97.3 2.2 51.6 47.7 0.4 27.6Cu20Mo30Ni50 60 days MoNi/fcc1/fcc2 49.1 50.5 0.8 45.9 14.1 72.4Cu80Mo10Ni10 60 days bcc/fcc 96.7 1.4 0.1 11.5Cu32Mo5Ni63 45 days fcc1/fcc2 2.1 51.1 6.0 67.3Cu30.7Mo6.3Ni63 60 days fcc1/fcc2 1.6 48.1 8.2 68.9

1000 Cu10Mo25Ni65 45 days MoNi/fcc 49.4 49.7 15.0 71.4Cu20Mo60Ni20 45 days bcc/MoNi/fcc 98.0 1.7 51.0 47.6 1.5 31.8Cu10Mo60Ni30 45 days bcc/MoNi/fcc 97.8 1.9 52.3 46.7 1.0 31.1Cu20Mo30Ni50 45 days MoNi/fcc1/fcc2 47.4 50.4 1.5 42.5 14.4 70.8Cu20Mo20Ni60 45 days MoNi/fcc1/fcc2 47.4 50.8 1.5 42.5 14.2 70.5Cu30.7Mo6.3Ni63 45 days fcc1/fcc2 2.5 50.0 8.8 66.6

1100 Cu10Mo25Ni65 30 days MoNi/fcc 51.3 48.3 16.3 68.2Cu20Mo60Ni20 30 days bcc/MoNi/fcc 98.0 1.8 51.3 47.4 1.3 34.1Cu10Mo60Ni30 30 days bcc/MoNi/fcc 97.7 1.9 51.0 47.7 1.5 34.0Cu20Mo30Ni50 30 days MoNi/fcc1/fcc2 49.3 49.0 3.3 43.8 14.5 66.6Cu20Mo20Ni60 30 days MoNi/fcc1/fcc2 50.3 48.5 2.9 43.5 14.9 66.6Cu80Mo10Ni10 30 days bcc/fcc 98.6 0.4 0.2 15.0Cu70Mo10Ni20 30 days bcc/fcc 98.2 0.4 0.4 24.6

1200 Cu10Mo25Ni65 20 days MoNi/fcc 47.5 50.1 20.1 70.0Cu20Mo60Ni20 20 days bcc/MoNi/fcc 97.3 1.9 47.8 47.9 2.4 36.2Cu20Mo30Ni50 20 days MoNi/fcc1/fcc2 47.8 47.9 2.5 39.7 18.3 65.4Cu80Mo10Ni10 15 days bcc/liquid 98.3 1.0 0.2 13.8Cu70Mo10Ni20 15 days bcc/liquid/fcc 96.4 1.4 1.0 20.5 2.4 32.4

1300 Cu10Mo25Ni65 15 days MoNi/fcc 47.7 50.5 23.2 68.1Cu20Mo60Ni20 10 days bcc/liquid/fcc 95.4 3.2 2.4 27.6 21.6 63.8Cu20Mo30Ni50 10 days bcc/liquid/fccCu80Mo10Ni10 7 days bcc/liquidCu70Mo10Ni20 7 days bcc/liquid

Fig. 3. X-ray diffraction patterns obtained from (a) Cu30.7Mo6.3Ni63 (at.%) and (b)Cu20Mo20Ni60 (at.%) alloys annealed at 1000 ◦C for 45 days.

97.1 2.2 2.7 27.9 21.4 63.898.7 0.9 0.7 14.598.3 0.8 0.5 24.2

confirmed and well marked by different symbols. The above exper-imental results suggest the existence of the miscibility gap of thefcc phase in the Cu–Mo–Ni ternary at 1000 ◦C, while there is only adetermined fcc single-phase region in the Kozlova’s previous work,as shown in Fig. 1. It should be noted that the limited examination(only by the metallographic observation) of the insufficient alloysamples may result in the doubtful conclusion. A three-phase equi-librium (bcc + fcc + MoNi) was identified in the Cu20Mo60Ni20 (at.%)alloy annealed at 1200 ◦C, and indicated in Fig. 2(c), where the fccphase is completely wrapped by the MoNi phase. In addition, thetwo-phase microstructure (MoNi + fcc) in the Cu10Mo25Ni65 (at.%)alloy annealed at 1300 ◦C for 15 days was shown in Fig. 2(d), wherethe MoNi phase with the flat plate shape distributes in the matrixof the fcc phase.

All the equilibrium compositions of the Cu–Mo–Ni ternary sys-tem at 900 ◦C, 1000 ◦C, 1100 ◦C, 1200 ◦C and 1300 ◦C determined byEPMA are listed in Table 1, which shows that (a) the miscibility gapof the fcc phase exists at 900 ◦C, 1000 ◦C, 1100 ◦C and 1200 ◦C, anddisappears at 1300 ◦C, (b) the solubility of Cu in the MoNi phaseat 900 ◦C, 1000 ◦C, 1100 ◦C and 1200 ◦C was measured to be about0.5 at.%, 1.5 at.%, 1.7 at.% and 4.0 at.%, respectively, (c) the elementsCu and Ni are almost insoluble in the bcc phase.

It is interesting to find that the as-cast Cu20Mo20Ni60 (at.%),Cu20Mo30Ni50 (at.%), Cu10Mo60Ni30 (at.%), Cu70Mo10Ni60 (at.%),Cu20Mo60Ni20 (at.%) and Cu80Mo10Ni10 (at.%) alloys appear theseparated macroscopic morphologies, which are caused by the liq-uid phase separation on cooling, while the as-cast Cu10Mo25Ni65

(at.%), Cu32Mo5Ni63 (at.%) and Cu30.7Mo6.3Ni63 (at.%) alloysappear the homogenous microscopic morphologies. Typical as-castmicrostructures are shown in Fig. 4(a) and (b). Such microstructuresare similar to the ones in the Cu–Fe, Cu–Cr and Cu–Co base alloys,as reported in our previous works [16–18].

40 Y. Yu et al. / Materials Chemistry and Physics 125 (2011) 37–45

Fa

3

C[p

G

wsi

w

a

Cics

Table 2The thermodynamic parameters used in the present calculation.

Thermodynamic parameters in each phase (J mol−1) Reference

Liquid phase : (Cu, Mo, Ni)10Lliq

Cu,Mo = 57285.4 + 2.0 T [23]

1LliqCu,Mo = −1200 [23]

0LliqCu,Ni

= 12048.61 + 1.29893 T [24]

1LliqCu,Ni

= −1861.61 + .94201 T [24]

0LliqMo,Ni

= −46540 + 19.53 T [22]

1LliqMo,Ni

= 2915 [22]

0LliqCu,Mo,Ni

= 100,000 This work

1LliqCu,Mo,Ni

= 150,000 This work

2LliqCu,Mo,Ni

= −70,000 This work

Fcc phase (A1): (Cu, Mo, Ni)10Lfcc

Cu,Mo = 83144 [23]0Lfcc

Cu,Ni= 8047.72 + 3.42217 T [24]

1LfccCu,Ni

= −2041.3 + 0.99714 T [24]0T fcc

c = −935.5 [24]1T fcc

c = −594.9 [24]0ˇfcc = − 0.7316 [24]1ˇfcc = − 0.3174 [24]0Lfcc

Mo,Ni= 4803.7 − 5.96 T [22]

1LfccMo,Ni

= 10, 880 [22]0Lfcc

Cu,Mo,Ni= −95190 + 30 T This work

1LfccCu,Mo,Ni

= −30, 000 This work2Lfcc

Cu,Mo,Ni= −33, 000 This work

Bcc phase (A2) : (Cu, Mo, Ni)10Lbcc

Cu,Mo = 82, 313 [23]0Lbcc

Cu,Ni= 8047.72 + 3.42217 T [24]

1LbccCu,Ni

= −2041.3 + 0.99714 T [24]0Lbcc

Mo,Ni= 46, 422 [22]

0LbccCu,Mo,Ni

= 1LbccCu,Mo,Ni

= 2LbccCu,Mo,Ni

= 0 This work

Ni3Mo phase : (Ni)3(Mo)10GNi3Mo

Ni:Mo− 0Gbcc

Mo − 30GfccNi

= −4199 − 7.00 T [22]

Ni4Mo phase : (Ni)4(Mo)10GNi4Mo

Ni:Mo− 0Gbcc

Mo − 40GfccNi

= −4330 − 9.21 T [22]

NiMo phase : (Ni)24(Mo,Ni)20(Cu,Mo)120GNiMo

Ni:Ni:Mo− 120Gbcc

Mo − 440GfccNi

= −1030 − 93.5 T + 13.5 T ∗ ln(T) [22]0GNiMo

Ni:Mo:Mo− 320Gbcc

Mo − 240GfccNi

= −212, 100 + 1089 T − 142 T ∗ ln(T) [22]0GNiMo − 200Gbcc − 240Gfcc − 120Gfcc = 170, 000 This work

ˇ = xiˇ + xixj ˇ (xi − xj) , (5)

ig. 4. (a) Homogeneous microscopic morphology of the Cu10Mo25Ni65 (at.%) alloynd (b) separated microscopic morphology of the Cu20Mo20Ni60 (at.%) alloy.

.2. Thermodynamic modeling

The Gibbs free energies of the liquid, fcc and bcc phases in theu–Mo–Ni system are described by the sub-regular solution model19]. According to this model, the molar Gibbs free energy of �hase in the Cu–Mo–Ni system is given by:

�m =

∑i=Cu,Mo,Ni

x0i G�

i+ RT

∑i=Cu,Mo,Ni

xi ln x�i

+ exG� + mgG�, (1)

here G�i

is the molar Gibbs free energy of pure element i in thetructure � phase, and the term exG� is the excess free energy, whichs expressed by the Redlich–Kister polynomials [18] as:

exG� = xCuxMoL�Cu,Mo + xCuxNiL

�Cu,Ni + xMoxNiL

�Mo,Ni

+xCuxMoxNiL�Cu,Mo,Ni,

L�i,j

=n∑

m=0

mL�i,j

(xi − xj)m,

L�Cu,Mo,Ni = xCu

0L�Cu,Mo,Ni + xMo

1L�Cu,Mo,Ni + xNi

2L�Cu,Mo,Ni,

(2)

here L�i,j

is the interaction parameter in the i − j binary system,

nd the L�Cu,Mo,Ni

corresponds to the interaction parameters in the

u–Mo–Ni ternary system. The coefficient of nL�Cu,Mo,Ni

is evaluatedn the present work. In the Cu–Mo–Ni ternary system, the magneticontribution to Gibbs free energy for the fcc and bcc phases is con-idered because magnetic transformation occurs in these phases.

Ni:Mo:Cu Mo Ni Cu0GNiMo

Ni:Ni:Cu− 440Gfcc

Ni− 120Gfcc

Cu = 400, 000 This work0LNiMo

Ni:Ni:Cu,Mo= 50,000 This work

The term mgG� is expressed as follows [20]:

mgG� = RT ln(ˇ� + 1)f (��), (3)

where ˇ� is a quantity related to the total magnetic entropy, whichin most cases is set to Bohr magnetic moment per mole of atoms; ��

is defined as T/T�c , and T�

c is the critical temperature for magneticordering, T�

c and ˇ� are described by the following expression:

T�c =

∑i

xiT�i

+ xixj

n∑m=0

mT�i,j

(xi − xj)m, (4)

�∑

�n∑

m � m

i

i

m=0i,j

where T�i

is the curie temperature for pure component, and the mT�i,j

is the interaction term between the elements i and j (i, j = Cu, Mo

Y. Yu et al. / Materials Chemistry and Physics 125 (2011) 37–45 41

Mo–N

oma

f

f

wt

tMmaitMiCe

Fig. 5. Calculated binary phase diagrams of the (a)

r Ni and i /= j). The ˇ�i

is the Bohr magnetic moment for pure ele-ents, and the f(��) represents the polynomials obtained by Hillert

nd Jarl as follows [21]:

(��) = 1 − 1D

[79�−1

140P+ 474

497

(1P

− 1

)(�3

6+ �9

135+ �15

600

)]for � ≤ 1, (6)

(��) = − 1D

(�−5

10+ �−15

315+ �−25

1500

)for � > 1, (7)

here D = (518/1125) + (11692/15975)((1/P) − 1) and P depends onhe structure, 0.4 for the bcc structure and 0.28 for the others.

No ternary intermediate phase was found in the Cu–Mo–Niernary system. There are three compound phases namely MoNi4,

oNi3 and MoNi in the Mo–Ni binary system. Because no experi-ental data on the solubility of Cu in the MoNi4 and MoNi3 phases

re available, the thermodynamic models of these two phasesn the Cu–Mo–Ni ternary system are treated as binary phases in

he previous literature [22]. However, the solubility of Cu in the

oNi phase cannot be ignored on the basis of the experimentalnformation obtained in the present work. The MoNi phase in theu–Mo–Ni ternary system is formulated using the sublattice model,.g. (Ni)24(Mo,Ni)20(Cu, Mo)12, and the molar Gibbs free energy of

i [22], (b) Cu–Mo [23] and (c) Cu–Ni [24] systems.

the MoNi phase in the Cu–Mo–Ni ternary system can be expressedas follows:

GMoNim =

∑i

∑j

yIIi yIII

j GNi:i:j

(i=Mo,Ni; j=Cu,Mo)

+ 20RT∑

i

yIIi ln yII

i + 12RT∑

j

yIIIj ln yIII

j

+yIIMoyII

NiyIIIj

(∑nLNi:Mo,Ni:i

(yII

Mo − yIINi

)n)

(j=Cu,Mo)

+yIIi yII

CuyIIIMo

(∑nLNi:i:Cu,Mo

(yII

Mo − yIINi

)n)

(i=Mo,Ni)

, (8)

where GMoNiNi : i : j is the Gibbs free energy of the MoNi phase,when the first sublattice is occupied by the element Ni, the secondone is occupied by the element i (i = Mo, Ni), and the third one isoccupied by the element j (j = Cu, Mo). LNi:Mo,Ni:j is the interaction

energy between Mo and Ni in the second sublattice, when the firstone is occupied by the element Ni, and the third one is occupiedby the element j (j = Cu, Mo).The LNi:i:Cu,Mo is the interaction energybetween Cu and Mo in the third sublattice, when the second one isoccupied by the element i (i = Mo or Ni), and the first one is occupied

42 Y. Yu et al. / Materials Chemistry and Physics 125 (2011) 37–45

) 900

bo

3

sbse

Fig. 6. Calculated isothermal sections of the Cu–Mo–Ni system at (a

y the element Ni. These parameters of interaction energies wereptimized in the present work.

.3. Optimization and calculation of phase equilibria

In the Cu–Mo–Ni ternary system, all phase diagrams in the threeubbinary systems, Mo–Ni [22], Cu–Mo [23] and Cu–Ni [24] haveeen thermodynamically assessed in the previous literature, ashown in Fig. 5. In the present work, the thermodynamic param-ters of the three binary systems were directly adopted from the

◦C, (b) 1000 ◦C, (c) 1100 ◦C, (d) 1200 ◦C, (e) 1300 ◦C, and (f) 1400 ◦C.

previous assessments. Because there is no experimental data onthe thermodynamic properties of the Cu–Mo–Ni ternary system areavailable, the ternary thermodynamic parameters were optimizedbased on the experimental data of the phase equilibria determinedand solidified microstructures investigated in this work. All the

thermodynamic parameters optimized in the present work andprevious studies are listed in Table 2.

The calculated isothermal sections at 900 ◦C, 1000 ◦C, 1100 ◦C,1200 ◦C, 1300 ◦C, 1400 ◦C and 1500 ◦C, compared with the exper-imental results determined in the present work, are shown in

Y. Yu et al. / Materials Chemistry and Physics 125 (2011) 37–45 43

F i(at.%)

Fab(isTpa(f(t1m3famtuosa

ig. 7. Calculated (a) isothermal section at 300 ◦C, (b) vertical section at Cu (at.%): N

ig. 6(a)–(f), respectively, where the calculated results are in goodgreement with the experimental data marked by different sym-ols. As can be seen in Fig. 6(a), there are two three-phase regionsbcc + MoNi + fcc and MoNi + fcc1 + fcc2) and a fcc-type miscibil-ty gap (fcc1 + fcc2) existing in the Cu–Ni-rich side, which is theame to the type of phase relationship shown in Fig. 6(b)–(d).he liquid phase is in equilibrium with the bcc phase and fcchase in the isothermal sections at 1100 ◦C (Fig. 6(c)), respectively,nd a three-phase region (L + bcc + fcc) also exist in Fig. 6(d) ande). In the isothermal section at 1300 ◦C shown in Fig. 6(e), thecc-type miscibility gap (fcc1 + fcc2) and the three-phase regionMoNi + fcc1 + fcc2) all disappear, which indicates that the criticalemperature of the fcc-type miscibility gap is between 1200 and300 ◦C. In order to systemically study the behavior of fcc-typeiscibility gap, the isothermal section of the Cu–Mo–Ni system at

00 ◦C, vertical section at Cu (at.%): Ni (at.%) = 1:1, and metastablecc-type miscibility gap in the Cu–Mo–Ni system were calculatednd shown in Fig. 7(a) and(c), respectively. In Fig. 7(a), the fcc-typeiscibility gap completely extends from the Cu–Ni side. Then, with

he increase of the temperature, the fcc-type miscibility gap grad-ally deviates from the Cu–Ni side on the basis of the presentlybtained experimental information presented in Fig. 6(a)–(d). Thispecial phenomenon is considered to be related to the effect of Modditions to the fcc-type miscibility gap in the Cu–Mo–Ni system.

= 1:1, and (c) metastable fcc-type miscibility gap in the Cu–Mo–Ni ternary system.

The calculated vertical section at Cu (at.%): Ni (at.%) = 1:1 of theCu–Mo–Ni system (Fig. 7(b)) obviously show that the fcc-type mis-cibility gap (fcc1 + fcc2) exists away from the Cu–Ni side at hightemperatures, which evidently indicates that Mo additions can sig-nificantly stabilize the fcc-type miscibility gap in the Cu–Mo–Nisystem at high temperatures, and is finally proved according tothe calculated results of the metastable fcc-type miscibility gap inthe Cu–Mo–Ni ternary system (Fig. 7(c)). Fig. 8(a)–(d) show thecalculated vertical sections at 50 at.% Ni, 60 at.% Ni, 63 at.% Ni and65 at.% Ni, respectively, and the calculated results are in agree-ment with the experimental data. Moreover, by comparing withthe vertical sections shown in Fig. 8(a)–(d), it is clearly found thatNi additions can significantly decrease the critical temperature ofliquid miscibility gap and reduce the region of the liquid miscibil-ity gap (L1 + L2). Fig. 9 shows the calculated liquidus projection inthe Cu–Mo–Ni system, where the alloys with the separated andhomogenous microscopic morphologies are marked by the sym-bols of squares and triangles, respectively. As can be obviouslyseen in Fig. 9, the miscibility gap in the liquid phase extends from

the Cu–Mo side to the Ni-rich corner, and the calculated criticaltemperature of the liquid miscibility gap is 1337.6 ◦C at the com-position point (4.1 at.% Mo and 58.1 at.% Ni) in Fig. 9. The presentlycalculated liquidus projection has two main differences from thatreported by Gupta [10]: the type of invariant reaction related to

44 Y. Yu et al. / Materials Chemistry and Physics 125 (2011) 37–45

Fig. 8. Calculated vertical sections of the Cu–Mo–Ni system at (a) 50 at.% Ni, (b) 60 at.% Ni, (c) 63 at.% Ni, and (d) 65 at.% Ni.

Fig. 9. Calculated liquidus projection in the Cu–Mo–Ni system compared to the macroscopic morphologies of the as-cast alloys. The symbols represent separated (�) andhomogeneous (�) macroscopic morphologies.

Y. Yu et al. / Materials Chemistry an

Table 3Calculated invariant reactions in the Cu–Mo–Ni ternary system.

Invariant reaction Type Temperature (◦C) Composition (at.%)

Cu Mo Ni

teMasatluTT

4

(

(

(

[

[

[

[

[

[

[

[

[[

L1 ↔ L2 + fcc P 1337.6 37.8 4.1 58.1L1 + L2 ↔ bcc + fcc E′ 1323.2 9.3 27.0 63.7L + bcc ↔ MoNi + fcc U 1319.3 7.2 29.0 63.8

he MoNi phase and the Ni-rich liquidus curve, which are consid-red to give explanation as follows: (1) the solubility of Cu in theoNi phase, which will affect or change the type of the invari-

nt reaction related to the MoNi phase, is measured to be muchmaller in this work than that in the reported liquidus projectionnd (2) Ni additions can significantly decrease the melting point ofernary Cu–Mo–Ni alloys, which results in the shape of the Ni-richiquidus curve towards Ni-rich corner. Therefore, the calculated liq-idus projection in the present work is reasonable and creditable.he details about the calculated invariant reactions are listed inable 3.

. Conclusions

1) The isothermal sections at 900 ◦C, 1000 ◦C, 1100 ◦C, 1200 ◦C and1300 ◦C in the Cu–Mo–Ni system were determined by met-allography, EPMA and XRD. The experimental results firstlyfound that the fcc-type miscibility gap exists at 900 ◦C, 1000 ◦C,1100 ◦C and 1200 ◦C, and disappears at 1300 ◦C, and the solu-bility of Cu in the MoNi phase at 900 ◦C, 1000 ◦C, 1100 ◦C and1200 ◦C are about 0.5 at.%, 1.5 at.%, 1.7 at.% and 4.0 at.%, respec-tively.

2) The as-cast Cu20Mo20Ni60 (at.%), Cu20Mo30Ni50 (at.%),Cu10Mo60Ni30 (at.%), Cu70Mo10Ni60 (at.%), Cu20Mo60Ni20(at.%) and Cu80Mo10Ni10 (at.%) alloys appear the separatedmacroscopic morphologies, which are caused by the liquidphase separation on cooling, while the as-cast Cu10Mo25Ni65(at.%), Cu32Mo5Ni63 (at.%) and Cu30.7Mo6.3Ni63 (at.%) alloysshow the homogenous microscopic morphologies.

3) A consistent set of optimized thermodynamic parameters basedon the experimental data have been derived for describing theGibbs free energies of the solution phases and intermetalliccompounds in the Cu–Mo–Ni system, which leads good agree-ment between calculation and most of the experimental data.

[[[[

[

d Physics 125 (2011) 37–45 45

Acknowledgements

This work was supported by the Ministry of Education of China(Grant No. 707037) and the Ministry of Science and Technology ofChina (Grant Nos. 2009DFA52170 and 2009AA03Z101). The Sup-port from Fujian Provincial Department of Science & Technology(Grant No. 2009I0024), Xiamen City Department of Science & Tech-nology (Grant No. 3502Z20093001), and Aviation Science Fund(Grant No. 2009ZF68010) are also acknowledged. We also thankfor the support from the Global COE project (Japan).

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