Displacement reactions between metal and liquid sulfides

Materials Chemistry 6 (1981) 147 - 164

DISPLACEMENT REACTIONS BETWEEN METAL AND LIQUID SULFIDES

S.R. SHATYNSKI D e p a r t m e n t o f Materials Engineering - Rensselaer P o l y t e c h n i c I n s t i t u t e - T R O Y , N . Y . (U.S.A.)

Received 19 January 1981; accepted 11 March 1981 Abstract - Displacement reactions between liquid sulfides and metals are studied and compared with the resultant morphologies previously observed in solid-state sulfide-metal and oxide-metal displacement reactions. In all cases the aggregate morphology was observed indicating that the rate controlling step in a tentatively assumed layered arrangement would be diffusion of atomic sulfur through the product metal. The experimentally observed morphologies for Nb-Ni 3 $2 (1), Nb-FeS (1) and Cu-Ni3 S 2 reaction couples are reported and rationalized. Detailed analysis of the liquid sulfide displacement reactions considered are not possible because of the lack of available thermodynamic and diffusivity data for the sulfide systems.

INTRODUCTION In high temperature multiphase composites, the occurrence o f displacement reactions between metal matrices and compound fibers, particles or precipitates is an important consideration. A solid-state displacement reaction between a metal and an ionic compound results in the formation of two or more products. Consider the simple solid-state displacement reaction between a metal Me and a sulfide MxS: u Me + MxS ~ M e u S + xM.

(1) 0390-6035/811030147-1752.00/0 Copyr~ht © 1981 by C~q]SDRS .R.I.. All fights of reproduction in any form reserved

148 Such a reaction will proceed if the Gibb's energy change per mole of sulfur is

AG AG~%s =

-

(2)

aC~x s

where the four phases each exist in their standard states, which is pure solids. Equation (2) can be rewritten in terms of the sulfur activities: AG =

RT

In

2

Ps2 (Me/MevS)

(3)

Ps~(M/MxS)

where Ps2(Me/MevS) and Ps2(M/MxS ) are the sulfur partial pressures for the coexistence of Me and MevS and M and MxS respectively. Wagner I first considered the morphology and kinetics of displacement reactions. Two morphological arrangements were then proposed for the products i) a layered structure shown in Fig. 1 a and ii) an aggregate arrangement shown in Fig. lb.

Ks

- - metal formed

M --sulfide

formed

M%S -metal

consumed

(al

Me

MxS sulfide

M.7 [- __ metal

formed--

formed

M," MevS

metal

I"

consumed-- " e°

Me

(,b)

Fig. 1 The possible morphologies for displacement reactions in the solid state (from Wagner I ).

149 Rapp et al. 2"4 more recently, theoretically predicted the morphologies and reaction rates of metal/oxide and metal/sulfide displacement reactions using the pertinent thermodynamic and diffusion data. It was noted that the analysis also should apply to nitrides, carbides, silicides, borides and other compounds. The displacement reaction between a metal M and the lowest sulfide of M, MxS to yield the metal M and the lowest sulfide of Me, i.e., MevS are considered. The standard states for the components in Equations (2) and (3) can now be defined as: the metal saturated with respect to its lowest sulfide and the sulfide saturated with respect to the metal. It is assumed that i) there is negligible mutual solubility between M and Me, M and MevS, Me aiad MxS and MevS and MxS, ii) no other binary or ternary compounds are formed, iii) the ionic phases are electronic conductors and iv) the product sulfide exhibits only cation diffusion. Wagner s in a paper concerning the morphology of diffusion-controlled oxidation proposed a criterion for the stability of flat scale growth. This criterion was extended to displacement reactions by Rapp et al. 3. As shown in Fig. 2, let the M/MevS interface be wavy. Assume that local equilibrium is achieved at each interface and that

MxS

MevS

M e.='~ Me

Fig. 2 - The layered morphology for a displacement reaction with a tentatively assumed wavy MIMerS interface (from Rapp et al.3 ).

the reaction rate is diffusion limited. If the growth of MevS is limited by diffusion of Me÷z cations then the flux of Me÷z cations that arrives at position II exceeds the flux that arrives at I and a layered structure results. Moreover, if the growth rate of MevS is limited by the diffusion of S through the M phase then the flux at position II exceeds that of position I

150 and a serrated or aggregate structure results. Thus the morphology of the reaction products is determined by the rate controlling step in the growth of the MevS phase. To evaluate the criterion, assume only a small S activity gradient across the product metal M and then a large Gibbs energy change across the product sulfide MevS. Then the 'rate of growth of the MevS is equivalent to sulfidation of pure Me in a gas with a sulfur activity corresponding to the coexistence of M and MxS. This can thus be described by Wa~ner s parabolic oxidation rate theory ( A x ) 2 = 2 t 1---I Ps2 Zcat (D*at) dlnPs2 i t = 2 k p ( M e v S ) t t 2 idp~2 I Zanl t

(4)

tt

where Ps2 and Ps2 are the sulfur activities at the metal/sulfide and sulfide/gas interfaces, respectively, Ax is the thickness change of the MevS phase, Dca * t is the cation self-diffusivity in MevS , and Zca t and Zan are the cation and anion valences respectively, and Kp (MevS) is the parabolic rate constant for tile growth of MevS. By knowing the Ps~ dependence of Dc*at, one can then evaluate the kp (MevS). If, however, the entire Gibb's energy change is across the product metal M, then the rate of growth of M is: ..7- =~" m m m ACs dt VM Js (M)= n'VM Ds ~(

15)

where X is the thickness of M, "n"is the number of equivalents of M added to the product (M) layer, V~mi is • the molar volume of M, js(M) is the flux of S in the product metal M, D s i s the diffusivity of S in M, ACs is the difference in S content across the product M layer. From Sievert's law: 1/2 Cs = KM Ps2

-

NsM m VM

(6)

where KM is Sievert's law constant and NM is the atom fraction of S in M. Substituting Equation (6) into Equation (5): m DsM Ks (p,l/2 r 1/2 X2 =2 [n" - , VM ,-s: -Ps2 )]t=2kp(M)t.

(7)

By comparing the magnitudes of the normalized parabolic rate constants for the product metal M and the product sulfide MevS, the morphology of the reaction can be determined. If kp (M)> kp (MepS) then the layered arrangement is stable while if kp(MevS ) > kp (M) then the aggregate is stable. Upon comparing this analysis with experimental observations of oxides and sulfides, good agree-

151 ment was obtained. Tile aggregate morphology can be further classified into the lamellar and interwoven morphologies 2,4. The interwoven morphology is proposed to exist for reactions in which at least a limited solubility between the metals M and Me and tile sulfides MeuS and MxS exist. The resultant morphology is shown in Fig. 3a. If there is very limited solubility of Me in M and MeuS in Mx S then the lamellar aggregate is assumed to predominate. Tile lamellar aggregate morphology is shown in Fig. 3 b.

MxS

to)

Me M S x

~

M layer

Me

(bl

Fig. 3 - The morphology o f the two variations o f the aggregate morphology.

In the present study, this analysis is extended to displacement reactions between metals and liquid sulfides. A solid state displacement reaction is defined with the stipulation that the products formed must be solids. However, tfiere is no stipulation as to the state of the reactants. Reactions between Nb and liquid NiaS2, Nb and liquid FeS and finally one reaction where the product phase is liquid, Cu-NiaS2 reaction were studied.

152

EXPERIMENTAL PROCEDURE The experimental procedure is essentially the same as that used in previous studies of displacement reactions 2"4. Solid cylindrical specimens 1 cm in diameter and 0.06 cm thick of Cu > 99.9 pct where used for the reactions. The metals where stress relieved by annealing at approximately one-half of their melting point for five hours. The annealing was done in an argon atmosphere in the presence of titanium sponge which acted as an oxygen and nitrogen getter. Similarly sized-compacted powder discs of reagent grade NiaS 2 were made. The sulfide pellets were then sintered at half their melting point in an argon atmosphere for 48 hours. It was necessary to reduce all the nickel sulfides to its lowest sulfide Ni3S 2. Pellets were placed in a quartz capsule with metallic Ni. The capsule was flushed three times and back-filled with argon. The capsule was put into a furnace at a temperature of 600°C for one week. The total reduction to NiaS 2 was verified by X-ray analysis. All samples were polished to a flat surface and a mirror-like surface using 1 micron diamond polish. Samples of Cu and NiaS 2 were then placed in spring loaded contact and placed in a furnace for reactions. All reactions occurred in a flowing purified Ar atmosphere. The reactions of FeS and NiaS 2 with Nb were performed by milling a cylinder of Nb into a cup. The cup was f'flled with reactant sulfide powder and inserted into the furnace on an alumina boat. Another alumina boat containing a mixture of Nb and NbS powder was placed previously upstream. Purified Ar flowed over the reactants in a manner similar to the powder compact specimens. A reaction was terminated by removing the reaction couple to the cool end of the furnace and allowing it to cool in Ar. The flow rate was increased during this period to reduce oxidation by minimizing the leaking of air into the furnace. Samples were analyzed using standard optical microscopy. Tile samples were further examined with a scanning electron microscope (Materials Analysis Company Model 700) and an electron microprobe (Applied Research Laboratory).

RESULTS AND DISCUSSION Nb-Ni3S2 Reaction at 1000°C for 15 hours The reaction between solid Nb and liquid Ni3S 2 would result in the formation of solid product phases i.e., Ni and NbS. There is limited solubility between

153 Nb and Ni. From the phase diagram 7, only 6wt% Nb can dissolve in Ni while 2wt% Ni is soluble in Nb. The NbS phase has a hexagonal structure a and an ionic radius of .7A 9 while the solid NiaS 2 is tetragonal or fcc 7 and has an ionic radius of .78A 9. Therefore, one would expect the NbS and NiaS 2 solids to have limited solid solubility. It is assumed that this immiscibility can be extended to include liquid NiaS 2 and solid NbS. It has been noted that all observed sulfide displacement reactions produced an aggregate morphology. In light of this, the lamellar aggregate is the expected morphology based upon the solubility criterion 2"4 In the initial stages of the reaction, some Nb dissolution in the NiaS 2 would occur as shown in Fig. 4a. Because of the high diffusivity of Nb in liquid NiaS 2

Ni3S Ni_S

~

Nbfal

Ni nch NisS z Nb

NbS

Ni3S-

~

N

i

rich

NisSz

i~l~=lD'p

Nb

(c)

7 (e)

I¢ Nb

Fig. 4 - Proposed mechanism for the ]ormation o f the morphology for the N b - N i a S 2 (1) displacement reaction.

154 much of the liquid would saturate with Nb. Upon supersaturation, the precipitation of NbS at the Nb-Ni3S2 boundary would occur with the surrounding liquid becoming very rich in Ni as shown in Fig. 4b. Since Ni is solid at 1000°C, the viscosity of the Ni-r,.'ch Ni3S 2 would increase. Because of the convection, many of the smaller dendritic precipitates would break off, and because of the high viscosity would appear suspended in the liquid, this is schematically shown in Fig. 4c. The larger dendrites would remain attached to the Nb surface and thus would con. tinue to grow. Because of the hexagonal structure of NbS, the surface energy is

Fig. 5 - Scanning backscattered electron micrograph o f the morphology o] the displacement reaction between Nb and Ni3S2 at lO00°C for 15 hours (Mag. 50)0.

155 highly anisotropic. This anisotropy then promotes the dendritic growth to occur in only one direction and hence the lack of branching by the larger needles. Upon cooling there is a precipitation of elemental Ni from the supersaturated Ni-rich Ni3S 2. This is shown in Fig. 4d. This final microstructure is shown in Figs. 5 and 6. Fig. 5 illustrates the general morphology of the NbS dendrites. Notice the many small dendrites that appear in the Ni-Ni3S2 matrix. Fig. 6 shows a greater magnification of a typical dendrite. Note also the darker Ni precipitates in the matrix.

Fig. 6 - Scanning backscattered electron micrograph of product sulfide dendrite from the displacement reaction of Nb-NiaS2 at lO00°C for 15 hours (Mag. 660X).

Upon examination of the dendrite using microprobe scan, Figs. 7(a, b, c), the dendrite is rich in Nb and S while poor in Ni thus verifying the dendritic composition to be NbS while the matrix is Ni-NiaS2.

156

i

I

Fig. 7 - Distribution o f Nb, S and Ni for the displacement reaction between Nb and NiaS 2 at lO00°C for 15 hours.

The Nb-FeS Displacement Reaction at 1000°C for 24 hours There is very limited mutual solid solubility between Nb and Fe 7 . NbS has a hexagonal structure s and an ionic radius of 0.7A 9. FeS is also hexagonal i o and has an ionic radius o f 0.64,a, 9. Thus one would expect at least some solubility between FeS and NbS. Because of the very limited solubility of Nb and Fe, the lamellar aggregate morphology should result based upon the solubility criterion 2"4 for

157 displacement reaction morphology. Initially it is expected that the Nb near the phase boundary becomes rapidly saturated with Fe and S while the FeS(1) dissolves a considerable amount of Nb as shown in Fig. 8a. Upon supersaturation, the precipitation of NbS and Fe occurs as the phase boundary as shown in Fig. 8b. The region ahead of the reaction zone would not be supersaturated, due to the large solubility of NbS in FeS. Upon

NbS I

Fel

-

Nb

"

I

s !

.i.

I

'I ' FeS(ll

Nb,

,~,='z-fl I r . ~

,

I

,

,

(Nb, .RI) S smunon

I

t

! , FIS(I)

~'

I

I "

I

I

Fe

(Nb, FeIS

solution

tb)

{oi Fe

,, I

I

(Nb, Felg solution

'

tel Fig. 8 - The proposed model for the growth of the displacement reaction prod" ucts for the Nb-FeS(1) displacement reaction.

158 growth, the relative amount of each product is dependent upon molar volume considerations. No molar volume data is available for NbS. Continued growth results in a large region of Fe solid solution in Nb, adjacent to the reactant Nb phase, and the two phase region of aggregate morphology, Fig. 8 c. Ahead of this two phase zone is a region of (Nb, Fe) S liquid. Fig. 9 illustrates the general morphology for the displacement reaction between Nb and liquid FeS at 1000°C for 24 hours. Examination of the reactant

Fig. 9 - Scanning backacattered electron micrograph o f the Nb-FeS(I) displacement reaction at lO00*C for 24 hours {Mag. 722X).

metal/product interface reveals the existence of a single phase of Fe and S in elemental Nb. In addition there is product on both sides of the crack, indicating the growth of this phase into the reactant Nb. Upon examination of the two phase zone in Fig. 10, the aggregate structure is evident. A distinctively different morphology from the Nb-Ni3S2 has resulted.

159

Fig. 10 - Optical micrograph o f reaction zone o f the Nb-FeS displacement reaction at I O00°C for 24 hours.

Cu-Ni3S 2 Displacement Reaction at 750°C for 47 hours Although this couple does not obey the criterion for solid state displacement reaction, nevertheless this couple will be evaluated in light of previous displacement reactions. Copper and nickel are completely miscible at 700°C 7. The solid solution between Ni3S 2 and Cu2S, however, is very limited 11 . KuUerud 11 reports that no ternary phases exist. He does report, however, the existence of a low melting eutectic at approximately 580°C and with an approximate composition of 12wt% Cu, 22wt% S and 66wt% Ni. Since Ni3S 2 and Cu2S are miscible liquids 11 , one would expect the formation of an interwoven aggregate structure based upon the solubility criterion for prediction of aggregate morphologies 2"4. The reaction begins by the dissolution of S and Ni in Cu and the solution of Cu in Ni3S 2. Because of the limited solubility between Cu2S and Ni3S2, dissolu-

160

tion of Cu in N i a S 2 should occur only in the immediate boundary region as seen in Fig. 1 la. Upon reaching the eutectic composition, localized melting occurs, as shown in Fig. 1 lb. The diffusion of Cu is enhanced inthe molten zone, where, upon supersaturation and precipitation of the product Cu2S and Ni-Cu alloy occurs, as shown in Fig. l i b . At the same time, more of the reactant sulfide rich in Cu melts at the eutectic composition, therefore the approximate width of the liquid phase remains constant. Because of the enhanced diffusivity of Cu cations and S in the molten zone, growth of the product zone should be limited by the dif-

eutee tic

liquid

Cu !

I

o

:

,

o

I

I

!

(m)

(bl Hi

OvsS Cu

NiB|2 t

t

Cu

t $ o e,--q

•

NIB| |

'

I # 0

NI

e

eutee~e Iiqul41

eu~Ne

Ilwid (¢)

(d)

Fig. 11 - Proposed model for the growth and morphology o f the Cu-Ni3S 2 displacement reaction at 750*Clot 47 hours.

161 fusivity of Cu cations in CuzS. The diffusivity of Cu in Cu2S was reported to be approximately 10 "s cm2/s. 12-14. Because of the large diffusivities, the reaction zone should appear very large. Also due to the large differences in molar volumes (2 V rn Cu2S = 56.4 cm a and 3 V nM i - 19.77 cm 3), the Cu2S should outgrow the product metal (Ni-Cu alloy). Therefore, precipitation of the product metal should occur only at the interface or behind it as shown in Fig. 1 lc. Tile observed morphology is shown in Fig. 12. The Ni-Cu alloy is certainly more abundant near the Cu interface. In addition, Fig. 12 illustrates the molar volume effect o f the Cu2S being _

Fig. 12 - Morphology o f the displacement reaction between Cu and NiaS 2 for 69.5 hours at 750"C (Mag. 10020.

162

Fig. 13 - Scanning backscattered electron micrograph o f product zone for the Cu-NiaS2 displacement reaction at 750"C for 69.5 hours (Mag. 2500X). much more abundant than the product Cu-Ni alloy. Voids in the product sulfide are polishing artifacts. This reaction couple is an example o f an interwoven aggregate morphology as demonstrated in Fig. 13.

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2. 3. 4. 5. 6. 7.

8.

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