Phase equilibria in ternary alloys based on iron-group metals and containing refractory metals (Mo, W, Nb, Ta)

Journal of the Less-Common

Metals, 144 (1988)

PHASE EQUILIBRIA IN TERNARY ALLOYS METALS AND CONTAINING REFRACTORY N. ASRAR,

L. L. MESHKOV

41

41 * 52

BASED ON IRON-GROUP METALS (MO, W, Nb, Ta)

and E. M. SOKOLOVSKAYA

Faculty of Chemistry, Moscow State University named after M. V. Lomonosov, Moscow-l 17234 (U.S.S.R) (Received

January

9, 1988;

in revised form

April 6, 1988)

Summary The phase equilibria in the Fe-W-Ta, Co-W-Ta, and Ni-W-Ta systems have been studied at 1273 and 1173 K, using physicochemical methods of analysis. Taking into consideration reported thermodynamic values of the intermetallic compounds formed in the analogous systems, their crystal structures and the nature of interaction between them, the importance of the near-neighbour diagram in assessing qualitatively crystallographic and thermodynamic factors, has been shown.

1. Introduction Modern industrial and technological developments demand new materials to work in aggressive atmospheres such as the combustion products of fuel and air, high-temperature catalytic reactors etc. [l]. In order to work satisfactorily in such environments, alloys must possess properties such as creep resistance, fatigue resistance and adequate resistance to hot corrosion [2]. The unique set of these required properties is obtained by using multicomponent alloys based on iron group metals and hardened by solutes and precipitates. Three important parameters which improve creep resistance of metals are low diffusivity, low stacking-fault energy and high elastic modulus [ 31. It has been shown [4] that the higher the melting temperature of the solute, the lower the diffusivity. Thus the higher melting solutes are to be preferred for better creep resistance. Decreases in the stacking-fault energy of metals of the iron group, due to a solute, are generally greater when the separation in the periodic table between the iron-group metal and the solute is greater [5, 61. Group VI elements, e.g. MO, W, have high melting points, large atomic diameters and high hardening coefficients. Similarly elements of group V, e.g. Ta, Nb, have high melting points and large atomic diameters and are well separated from the iron group metals. Tantalum and niobium are used as the elements to increase oxidation and hot-corrosion resistance [ 71. Use of these elements may, therefore, provide a unique solution of many 0022-5088/88/$3.50

@ Elsevier

Sequoia/Printed

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42

mechanical and physical problems of the alloys. As the theory of alloys has not yet been sufficiently developed to predict more complicated phase diagrams and optimum alloy compositions, a systematic study of the phase diagrams of these metal systems is required. In the present work three un-investigated ternary systems, Fe-W-Ta, Co-W-Ta and Ni-W-Ta, have been studied to improve our understanding of the behaviour of refractory metals of groups V and VI in iron, cobalt and nickel matrices. From these investigations, observations concerning intermetallic compound equilibria in such systems are reported here. Taking into account crystallographic and thermodynamic aspects of some analogous ternary systems and the tendency of complete series of solid solutions to form between isostructural phases, an attempt has been made to explain the structure of these phase diagrams. On the basis of this explanation it seems to be possible to predict phase equilibria in more complex un-investigated systems using the minimum possible experimentation.

2. Experimental details Ternary alloys of Fe-W-Ta, Co-W-Ta and Ni-W-Ta systems were studied in cast, homogenized and annealed states. Specimen alloys were prepared using the powder metallurgy technique. Compacted pellets of properly mixed powders (purity, 99.98%) were sintered in a vacuum furnace at a temperature of 1273 K for 24 h followed by a melting process in an argon-atmosphere arc furnace at 90 V. Heat treatment for the purpose of homogenization and the attainment of equilibrium, was carried out at 1573 K for 100 h in a furnace with a vacuum of 10P4 atm. Then for a long annealing, sample alloys were sealed in vacuum quartz capsules containing chips of zirconium for the preferential oxidation. Fe-W-Ta and Co-W-Ta ternary alloys were annealed at 1273 K and Ni-W-Ta alloys at 1173 K for 2000 h, followed by iced water quenching. Examination of these alloy samples involved optical metallography, X-ray diffractometry, hardness and microhardness tests and electron microprobe analysis. 3. Results 3.1. Fe- W-Ta system (Fig. 1) In this ternary system all the intermediate phases present in the Fe-W and Fe-Ta binary systems [8, 91, i.e. the Laves phases FezW and Fe,Ta and the p phases Fe,W, and FeTa, have been found. The isostructural 1-1phases, Fe,W, and Fe,Tas, appear in the ternary system up to a limit of 6 at.% tantalum and 8 at.% tungsten, and do not form a complete series of solid solutions, while the Laves phases form a complete series of solid solutions. This has been confirmed by X-ray diffractometry and microhardness tests (Fig. 2).

43

Fig. I, Ternary phase diagram of the Fe-W-Ta

system at 1173 K.

3.2. Co- W-Ta system [l O] (Fig. 3) All the binary phases of the Co-W and Co-Ta systems [ 11,121 exist in this ternary system. At 1273 K, CoSTa and Co,.,Ta dissolve 5 at.% tungsten, while CoTTab and CoTaa dissolve 9 at.% and CoaTa dissolves 13 at.%. The solubilities of tantalum in Co,W and Co,W, are 8 at.% and 18 at.% respectively. The P phases, Co,Ta, and C&W,, do not form any series of solid solutions. Like the Co-Mo-Ta system [13] here p phases also dissolve appreciable amounts of the constituents of the respective systems and exhibit different ternary equilibria. 3.3. Ni- W-Ta system [14] (Fig. 4) Keeping in view the dispute about the formation of the binary phases NiW and NIW:, in the Ni-W system, attention had been paid to analysing the binary and ternary alloys of Ni-W-Ta system {Fig. 41, which belong to the areas where formation of these phases has been reported previously tl.5 - 17 3. Metallography and X-ray phase analysis of these alloys, annealed at 1173 K for 2000 h, do not show the existence of the reported NiW and NiW2 binary phases. Figure 5 and Table 1 show that, at this temperature, Ni4W is found in equilibrium with the soIid solution of W-Ta. Microprobe results, showing the chemistry of the reaction zone of nickel-based matrix (alloyed with tungsten and ~nta~urn~ with tungsten fibre [US], do not give any information about the formation of these binary phases. Ni,W on the tungsten side and Ni,Ta, NizTa, NiTa and NiTa on the tantalum side, are all found and have the same crystal structure BSfound in

44 i c 1.84-

0.

780 _ -c

l.76-

A

A0

4.84

4.82

rn-

(a)

(b)

Fig. 2. (a) Crystal lattice parameters of the complete series of the solid solutions of ironbased Laves phases (f, reported values). (b) Microhardness of the complete series of solid solutions of iron-based Laves phases.

Fig. 3. Ternary phase diagram of the Co-W-Ta

system at 1273 K [lo].

45

ffi

_Wat% Fig. 4. Ternary

(b)

phase diagram

of the Ni-W-Ta

system

at 1173 K [la].

28

Fig. 5. Line representation of the X-ray diffractogram of the binary alloys SOat.%Ni50at.%W (a) and 36at.%Ni_64at.%W (b) showing the two-phase region of NieW (x) + LW-Ta (0).

46 TABLE

1

Results

of the X-ray

d(NiW)

(A)

36at.%Ni-64at.%W 2.27 2.11 2.08 2.01 1.96 1.93 1.82 1.72 1.64 1.53 1.42 1.22 1.19 1.18 1.09 1.05

SOat.%Ni-50at.%W 1.03 0.996 0.952 0.934 0.924 0.908 0.907 0.899 0.898 0.884 0.834 0.833 0.832 0.817

diffraction d(Ni4W)

2.887 2.070 2.020 1.807 1.771 1.679 1.625 1.506 1.446 1.430 1.347 1.331 1.291 1.278 1.266 1.159 1.122 1.088 1.073 1.037 1.017 0.981 0.946 0.909 0.905 0.903 0.900 0.885 0.879 0.830 0.797

patterns (A)

of the Ni-W

specimen

alloys

dexp (A)

1111

hkl

Intermediate phase

2.235 2.073

90 100

1.798 1.576 1.446 1.326 1.288 1.260 1.116 1.120 1.072 0.985 0.911 0.841 0.786

45 60 20 10 90 15 65 10 10 80 65 12 15

110 211 310 200 321 222 411,211 213 220 510 213 530 611,222 321 400

o1(W-Ta) Ni4W( A) A (Y A A

2.240

95 65 20 45 10 100 55 85 10 12 35 10 10

110 211 301 200 330 213 220 530 222,413 442 321,631 314 400

a

1.069 1.802 1.578 1.340 1.268 1.101 0.982 0.900 0.877 0.834 0.801 0.790

A, Q A CI A A A A, CY 01 (Y

A A (Y A A a! A A, QI A A, a A (Y

the binary system [19, 201. The solubility of tungsten in NisTa, Ni,Ta, NiTa and NiTa, and that of tantalum in Ni,W is not more than 5 at.%. All these phases have a small area of homogeneity of 4 at.% nickel. The solubility of iron, cobalt and nickel in the W-Ta solid solution does not exceed 1 at.%.

4. Discussion The investigated ternary systems exhibit interactions between the phases present in the constituent binary systems but formation of ternary

47

intermetallic compound does not take place. These results verify the prediction in ref. 13 and agree with the generalization given there about the construction of the phase diagrams being determined by complete dissolution of the constituent group V and VI refractory metals. Analysis of these results from that viewpoint shows that complete dissolution of the intermetallic compounds into each other takes place between the isostructural Laves phases of the iron-based system only. Complete series of the solid solutions between the Laves phases of the type Fe2M (M 5 group V - VI refractory metals) have also been found in the Fe-MO-W, Fe-Mo-Nb and Fe-W-Nb systems (Figs. 6 (i, iv, vii)). This indicates that complete dissolution of the isostructural Laves phases occurs in all the FeM(V)-M(V1) ternary systems and does not depend upon the diameter, electronic structure or electronegativity of the alloying atoms.

2

Nb

(ivl

M

l Vii)

W

Nb

Fig. 6. Ternary phase diagrams CO-W-MO at 1273 K [13], (iii) [22], (v) Co-Nb-Ta at 1373 K 1173 K [24], (viii) Co-Mo-Ta at

Nb

(ix)

w

of transition metals: (i) Fe-W-MO at 1173 K [13], (ii) Ni-Ta-Mo at 1173 K [21], (iv) Fe-Mo-Nb at 1173 K [13], (vi) Ni-MO-W at 973 K [23], (vii) Fe-Nb-W at 1273 K [13], (ix) Ni-Nb-W at 1173 K [25].

48

Complete series of solid solutions of equiatomic phases, i.e. p phases, form when the alloying metal atoms of the corresponding phases are isoelectronic, e.g. molybdenum and tungsten in the Fe-MO-W and CO-MO-W systems (Figs. 6 (i, ii)) and niobium and tantalum in the Co-Nb-Ta system (Fig. (6 v)). Interactions between the p phases do not occur in the Co-TaMO, Co-Ta-W and Ni-Mo-Ta (Figs. 6 (iii, viii) and Fig. 3) systems. The influence of the group factor on the appearance of complete series of the solid solutions between 1-1phases has not been observed in Ni-W-MO, Ni-W-Nb and Ni-W-Ta ternary systems (Figs. 6 (vi, ix) and Fig. 4) due to the absence of p phases in Ni-W binary systems of these ternary phase diagrams. Sokolovskaya et al. [26] have also challenged the formation of these phases in the Ni-W binary system. Recently ternary alloys in the Ni-W-Cr and Ni-W-B systems have been studied after very long heat treatments of 10574 - 11080 h at 1000 “C and 5550 - 27750 h at 950 ‘C, and no intermetallic compound except Ni,W has been found on the Ni-W side of these ternary systems [27, 281. It seems that these intermetallic compounds are metastable in nature and decompose after long annealing. It appears from the phase diagrams given in Figs. 1, 3,4 and 6 that the ternary phase equilibria in iron-based and nickel-based ternary systems are similar. The phase diagram of the cobalt-based systems, consisting of a comparatively large number of intermetallic compounds, is determined by the thermodynamic properties and occurrence of many phases having complex crystal structures (Table 2). Due to these reasons, it has not become possible to calculate or forecast the exact nature of the phase diagrams of such systems. Interactions of the iron-group metals with the groups V and VI refractory metals results in the formation of intermetallic compounds with different crystal structures and type of bonds. In the majority of cases, the physicochemical properties of even the closely related phases are quite different. These observations point towards the fact that in these intermetallic compounds appreciable restructuring in the electronic structure of the atoms takes place, which also causes changes in composition and conditions for equilibrium. X-ray diffractometry and physicochemical annalysis data given in Table 1, show that the majority of the intermetallic compounds which determine the structure of the phase diagrams have close-packed structures. In such types of intermetallic compounds, atoms are supposed to occupy maximum space in the crystal lattice in order to have minimum volume and maximum coordination number. Highly symmetric Laves phases with close-packed structure are formed by atoms of different diameters, rA > rz. Experimentally it has been proved that atoms interacting with each other in intermetallit compounds undergo a variation in size, no matter whether they were initially of larger or smaller size [38]. In explanation of this, Nevitt [39] proposed the existence of a strong interaction between the different atoms in the intermetallic compounds, which also reflects on the thermodynamic properties of these intermetallic compounds. Taking into account the fact that atomic size factor plays a deciding role in the stability of close-packed

Fe-/W6 CoNb Co7Tae Co,Moe Co,Mos cosw co*0 CosTa NiTa CoTaa

Coa.TTa FeNb FeTa Fe7M06

FeaZr FeaNb FeaTa FeaMo FeaW CoaNb CozTa CosNb

30 - 35 28 - 36 30 - 33 33.1 - 36 33.3 - 34 31 - 35 31 - 35 23 - 27 26.5 - 28 49.5 - 50.5 49 - 54 38.5 . 43 39 - 40 49 - 56 48 - 54 45 - 50 45 - 50 24 - 26 23.5 _ 25 25 66 - 67 66 - 67

( K)

Congruent Congruent Congruent Peritectoid Peritectoid Congruent Congruent Peritectoid Peritectoid Congruent Congruent Peritectoid Peritectoid Congruent Congruent Peritectoid Peritectoid Peritectoid Peritectoid Peritectoid Peritectoid Peritectoid

Tf.xm,tion

1948 1928 2048 1190 1323 1753 1903 1513 1793 1723 2123 1647 1910 1647 1923 1783 1963 1373 1293 1263 2060 2073

compounds

Homogeneity range (at.%)

of the intermetallic

Type of formation

properties

Physical

I~termetal~ic compounds

2

TABLE

7.058 4.827 4.81 4.745 4.727 6.762 6.745 4.74 4.75 4.92 4.93 4.754 4.746 4.95 4.928 4.767 4.761 5.13 5.13 5.17 6.216 6.12

a(A)

-

-

15.45 15.47 26.79 27.1 25.71 25.78 26.30 26.44 25.65 25.72 4.12 4.12 18.91 4.872 4.97

-

7.778 7.84 7.734 7.704

c(A)

metals

b(A)

Crystal lattice

of the transition

3.268 3.24 5.442 5.494 5.405 5.436 5.348 5.326 5.382 5.40 0.803 0.803 3.658 0.784 0.812

1.626 1.624 1.629 1.63 -

-

c/a

-@

(1300 K)

-

4.6 3.8 5.2 6.1

3.8 3.3

8.1 7.6 19.2 28.2 14.2 22.8 -

24.7 20.4 25.0

(kJ (g atom)-*)

-

-

-

0.6 0.4 1.8 2.1

0.4 0.9

7.2 4.6 2.1 4.4 3.9 2.1 0.2 1.2 -0.5 -

--AS (1300 K) (J k-l g atom)-‘)

11,29 27,30 9,3s. 32,33 8, 34 35,36 12,36 35,36 12,36 21 9 32,33 8, 34 35 12 33,37 11,33 11,33 33,37 12 20 12

References

50

crystal structures, it is essential to evaluate its actual contribution to the stability of the interacting phases in different phase diagrams. In order to gain some quantitative assessment of the influence of the geometrical principles of metal structure on the occurrence of phases and on their structural dimensions, a useful model has been proposed by Pearson [ 401. This model allows the atoms to be compressed until successively A-A, A-B and B-B contacts are formed. During a strong interaction the interatomic distance in the structure will be less than the corresponding distance in an ideal contact. This results in a change in the crystal lattice. A definite type of interatomic interaction in Laves phases can be assessed with the help of the near-neighbour diagram given in Fig. 7. The straight lines correspond to the ideal axial ratio (c/u = 1.63) and ideal atomic parameters of the Laves phases. Points represent deformation parameters, which have been calculated from X-ray phase analysis data of the intermetallic compounds. Despite the fact that absolute enlargement of the deformation parameters may depend on the selection of the values of DA and D, (DA = 2r, and DB = 2r,), the quantitative assessment of the influence of the interatomic contacts remains the same as the deformation parameters and radius ratio increase proportionally. In Fig. 7 all points lie within the A-B and B-B contact lines up to their crossing point. For all the Laves phases, points lie below the 12-6 A-B contact line. This means that the interatomic contact of the given type does not form in the structure of the intermetallic compound. Identical atoms (A-A) taking part in the formation of the Laves phase (e.g. Fe,Hf, FezZr etc.) will be compressed and (B-B) contacts will be slightly expanded in order to minimize the crystal lattice distortion of the intermetallic compound. In systems where the interatomic contact (A-A) lies below the contact line 4A-A, a complete series of solid solutions between Laves phases

0.15

-0.05

Fig. 7. Near-neighbour diagram for the Laves phases (MgZnz type) constructed for the ideal axial ratio c/a = 1.63. The coordination for the contact lines is indicated. Points represent binary phases.

51

forms (e.g. FezMo, Fe,W etc.). Analysis of the Laves phases reveals that enlargement of the periods of the crystal lattice is not much influenced by the compression of the large atoms during alloy formation. This is due to the difference in numbers of corresponding interatomic contacts, i.e. in these structures contacts (A-A), (B-B) and (A-B) are found in the ratio of 1:3:3. Undoubtedly, this model is simple but it provides a way of comparing the variation in the interatomic distances in the intermetallic compounds with those in the typical close-packed structures. On this basis it may be concluded that depending upon the composition, some definite types of atomic interactions influence the stability of the intermetallic compounds. The thermodynamic properties of the intermetallic compounds given in Table 2 show a correlation with their positions in the near-neighbour diagram. This indicates that the enthalpy of formation of intermetallic compounds decreases with an increase in the deformation parameter. Energy spent on the deformation and breaking of the contacts between the identical atoms A-A and B-B, is compensated by the formation of contact between atoms having different electronic configurations, i.e. A-B. The consistent proximity of the deformation parameter to the 6-12 A-B contact line confirms this fact. Intermetallic compounds Fez (MO, W) occupy positions below the 4A-A and 6-12 A-B contact lines and have relatively high enthalpies of formation. On the basis of X-ray phase analysis and thermodynamic studies, an evaluation of the role of some definite interactions of the atoms responsible for the stability of the intermetallic compounds can be made for the whole group of close-packed structures. In the case of intermetallic compounds having considerable area of homogeneity, i.e. p phases in the systems discussed above, each group of contacts will not be represented by a line but by a band. Thus, the positions of the points representing intermetallic compounds on the near-neighbour diagram permit size and thermodynamic factors to be assessed qualitatively together. Conclusions (i) The phase diagram of the iron-group metal with group V and VI refractory metals is characterized by the existence of complete solid solubility between two intermetallic compounds. (ii) Use of the Pearson’s model and reported thermodynamic properties of the intermetallic compounds has enabled us to assess the possible contribution of some definite interatomic interactions to the stability of the complex intermetallic compounds. References 1 M. J. Wahl, D. J. Maykuth and H. J. Hucek, “Hand book of Supemlloys”, Columbus, OH (1979) p. 1. 2 R. W. Fawley, “The Supemlloys”, Wiley, New York, 1972, p. 3.

Battelle,

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