The stability of intercalation compounds AgxTaSs

Solid State lonics 28-30 (1988) l 116-1122 North-Holland, Amsterdam

THE STABILITY OF INTERCALATION COMPOUNDS Ag,,Ta.S~ G.A. WIEGERS, A.G. GERARDS #, H. ROEDE, R.J HAANGE and B.A. BOUKAMP * Laborabory of lnorganic Chemistry, Materials Science Center of the University, Nyenborgh 16, 9747AG Groningen, The Netherlands Received 13 August 1987

The electrochemical properties of intercalates Ag~TaS2 prepared by high-temperature techniques were measured in a solid electrolyte (Agl) electrochemical cell. The EMF versus x curve at 444 K showed the presence of stage-2 and -1 phases for 0.22 < x < 0.35 and 0.55 < x < 0.68 respectively. From X-ray powder diffraction evidence was found that for x < 0.1 a disordered staged phase, presumably a third stage phase is stable. The EMF is small (less about 0.25 V) and the temperature coefficient of the EMF is positive. A simple model was used to explain the observed EMF versus x curve. The free enthalpy of intercalation of the stage- 1, stage-2 and stage-3 intercalates is expressed into ( 1 ) a binding energy of the silver atoms to the lattice, proportional to x; (2) a repulsion energy of silver atoms in a plane and between silver atoms in neighbouring planes; (3) an elastic energy needed to open a gap between sandwiches using the rigid plate model; (4) the configurational part of the entropy using the mean field approximation. Using the Einstein model of vibration for silver atoms one finds a vibrational part of the intercalation entropy AS,, which is proportional to .r, in the free enthalpy of intercalation the contribution TAS~ can be absorbed in ( 1 ). With five parameters the observed EMF versus x curve is reasonably well reproduced. One also finds that the difference in free enthalpy for n-staged phases with n > 2 and a two phase mixture of stage 2 and 2H-TaS2 is very small which explains why higher-staged phases are difficult to obtain.

1. Introduction The reaction of the [b metals copper and silver with transition metal dichalcogenides TX2 (T=Ti, Nb, Ta; X = S, Se) can 0e performed in an electrochemical cell even at room temperature which indicates that the mobility of silver and copper in these phases is relatively high. Structural studies have shown that the Ib metal atoms are in planes between TX2 which have about the same geometry as in the host TX2, the ionic conduction being two-dimensional. The phases MxTX2 (M = Ag, Cu) have electronic as well as ionic conduction and behave as solid-solu*.ion electrodes in an electrochemical cell. In conhast to the lithium intercalates which are single phase for 0 < x < 1 the Ib metal intercalates form several crystallographically distinct phases depending on composition and temperature. Present address: Billiton Research t3V, Arhnem, The Netherlands. " Present address: Technical University Twente, Enschede, The Netherlands.

~n this paper we describe the results of a study of the structures and thermodynamic properties of phases AgxTaS2 from measurements on powder compacts in an electrochemical cell. The kinetics of the intercalation process and the physical properties like the electrical transport and magnetic properties and ordering phenomena as studied by electron diffraction will be discussed in forthcoming papers.

2. Phase studies and structures The synthesis and structure of a stage-1 phase Ago.67TaSz was described first by van de Berg [1]. o ~ u~,tu~ai atuuic~ uy oL:uuzz a n u r r m u t

[.z ] oy A - r a y

powder techniques of phases AgxTaS2 prepared at high temperature showed that at room temperature three distinct phases exist: what was called a dilute first stage phase for 0 < x < 0.08 with all van der Waals gaps between sandwiches TaS2 filled with silver in a low concentration, a second stage phase (gal~s between sandwiches alternately empty and partially

0 167-2738/88/$ 03.50 O Elsevier Science Publishers B.V. ( North-Holland Physics Publishing Division )

G.A. Wiegers et aL/Intercalation compounds AgxTdS 2

filled with silver) for 0.17 < x < 0.38 and a first stage phase for 0.55
2H-TaS2 which therefore can be considered to be the host for intercalation. The sandwiches of the stageI and stage-2 phases are displaced with respect to the original stacking of sandwiches in 2H-TaS2 in the case that a gap is filled with silver, as is demonstrated in fig. 1; the stacking of sandwiches is the same as in 2H-MoS2. The silver atoms of the stage-1 and -2 phases are in tetrahedral coordination by sulfur atoms of neighoouring sandwiches. The tetrahedral sites form puckered honeycomb (hc) lattice planes as shown in fig. 2. Each hc lattice can be considered to consist of two interpenetreting sublattices a and [3. Per tantalum atom present there are two tetrahedral sites (one ct and one [3 site) available. Neighbouring sites are at a close distance (approximately 1 / 3 a x / ~ = 1.9 A) and for that reason preferentially not occupied simultaneously. Planes with silver atoms are as far apart as 13.2 A in the second stage phase and 7.2 A in the first stage phase. In a recent paper Scholz and Frindt [3 ] proposed the existence of two intermediate phases; we will discuss this in relation with the EMF versus composition curves in section 3. In our investigations the samples Ag,TaS2 were prepared from the elements at 900 °C during six days;

G/3

a

b

G

c

Fig. I. ( 1 1 20) sections of the structures of (a) 2H-TaS2 (b) second stage Ag,TaS> space group R3m, I/3 of the c-axis is shown (c) first stage Ag,TaS,, space group P63/mmc. The origins are displaced with respect to the conventional ones in order to demonstrate the shear transformations during intercalation.

I 1 17

\,/

I Fig. 2. The planes of silver between sandwiches TaS,; the lattice sites of the two interpenetrating subla~tices ~t(@ ) and 13(O ) form a honeycomb lattice: for the stage-1 and -2 intercalates the sites are both occupied.

Table 1 Crystallographic data of 2H-TaS2, stage-2 Ago sTaS2 and stage- l Ago.6TaS2.

2H-TaS2 stage-2 Ago sTaS2 stage-l Ago 6TaS2

a (A)

c (A )

space group

3.315 3.327( 1 ) ") 3.3329(5)

12.10 39.72(4) 14.439(3)

P63/mmc R3m P63/mmc

') Standard deviations in parenthcses.

the samples were slowly cooled to room temperature. The X-ray data of stage 2 and stage 1 (table 1 ) agree with those of Scholz and Frindt [2]. The powder patterns of Ago ~TaS2 and 2H-TaS2 (prepared a* 700°C) are almost the same. They show broad line; apart from reflections h k 0 and 0 0 l for 2H-TaS2 and reflections h k0 for Ago.~TaS2 which are sharp. The spacing of sandwiches is 6.05 A for 2H-TaS2 and 6.50 A for Ago.ITaS2 (the 0 0 2 reflection has a maximum intensity at about d=6.50 ~, with a shoulder at d= 6.05 A which indicates some 2H-TaS2 to be present). These observations indicate considerable stacking disorder in both substances (broad general reflections h k l). The occurence of broad lines also for the 0 01 reflections in the intercalate indicates that the spacing between sandwiches varies due to unequal densities Gf ~ilver in different layers, the extreme case being a disordered higher staged phase, presumably n=3. From the average distance between sandwiches TaS2 in 2H-TaS: ~,~ ,,.,,., ~ . . . . tim stage-2 and stage-i phase of 6.50 and 7.22 A respectively, one finds that for each occupied plane the increment in the distance between sandwiches is 1.10

1118

G.A. Wiegerset al./lntercalation compounds AgxTaS~

A. For a stage-3 compound one expects therefore an average distance of 6.43 A which shows that Ago~TaS., prepared at 900°C probably is a disordered stage-3 compound and not a dilute stage-I phase with all gaps occupied by silver in a low concentration. A sample Ago.~TaS2 annealed at 600 ° C for one week showed well separated reflections at d=6.05 A and dffi6.42 A corresponding to 0 0 2 of 2H-TaS2 and a basal reflection of presumably a disordered stage-3 phase with x about 0.15. High-temperature X-ray powder diffraction of Ago.3TaS2 using a Guinier-Simon camera (EnrafNonius) in the temperature range 300-1100 K did not show a transition to first stage as expected from thermodynamic considerations (increase of the configurational entropy). Such a stage transformation was for instance observed for stage-2 Nao 3TiSe2 [ 4 ]. The possible ordering of silver atoms on sites of the heneycomb (hc)lattice has been the subject of a number of experimental studies [ 5-8 ]. Ordering which occurs below about 300 K leads to supeflattices. Since this paper deals with measurements above 440 K the low-temperature ordering will not be discussed. Even in the basic lattice, ordering of silver is possible, viz. ordering of atoms and vacancies over sites of the two sublattices ~t and ~. From the phase diagram of the hc lattice by Sato and Kikuchi [ 9 ] using the pair approximation of the cluster variation method as well as from the renormalization group approach by Sabbaswamy and Mahan [ 10 ] it is clear that both sublattices in stage-1 and stage-2 AgxTaS~ are equally occupied by silver. A single crystal X-ray study of stage-2 Ago.z3TaS2 crystals prepared by vapour transport showed this indeed to be the case [ 11 ]. For Ago.6NbS2, isostructural with Ago.6TaS2the same was found [ 121.

tion of the applied pressure. The EMF measurements were performed in a three electrode cell of the type: Ag/AgI/Ag~TaS2/C with the Ag reference electrode placed in the solid electrolyte near the AgI/ AR~TaS2 interface. This cell is useful above 418 K, being the temperature of the ~ to a transformation cf Agl. Use was also made of the galvanostatic intermittent current titration technique (GITT) [ 13 ]. Details of the experimental techniques which were also used in our studies of phases AgxTiS2 and AgxNbS2 are given in our previous papers [ 14,15 ]. The EMF versus composition curve measured at 444 K is shown in fig. 3; the composition was changed coulometrically using a current of l0 -4 A/cm 2. At regular intervals the current was interrupted in order to measure the equilibrium voltage which was reached in a few minutes except for compositions x<0.1 where the kinetics became very slow. The mrve shows two-phase regions between 0.1 < x < 0.22 and 0.35
200

E (mY} 100

\ For electrochemical studies the samples ground to powder of 75 gm panicles were pressed at room tempe."mre into cylindrical or bar shaped tablets in a steel die. In most cases the tablets or pellets were sintered at 70 n "C for 2-3 h. Tablets with an area of 0.1 cm-" and a thickness of about 0.05 cm were used; cuts were made parallel and perpendicular to the direc-

0

, 0.1

I

0 I. ~.

,

013

~

X

~ 0.~.

,

, 0.5

~

q 0.6

Fig. 3. The compositional (x) dependence of the EMF of the electrochemical cell Ag[ Ag! [AgxTaS2 i C g40 K from measurements performed with samples with different composition, some ,,~,~ained bv coulnmetric titration.

G.A. Wiegers et aL/lntercalation compounds AgxTaS2

tion of silver ions in the stage-2 and -1 phases. No measurements were possible for x < 0.1 because of the very slow kinetics in that region. Our EMF versus composition curves may be compared with those from liquid electrolyte electrochemical cells at 300 K of Scholz and Frindt [ 3 ] using an electrochemical cell with a AgNO3 solution in glycerol and that of Left and Butz [ 16 ] using AgNO3 in acetonitrile. The curve of Scholz and Frindt [ 3 ] which was thought to represent equilibrium, was explained in terms of the phases already discussed and two other phases, viz. a stage-2a phase with Ag in octahedral holes and a stage-1 a phase with Ag in succesive planes in tetrahedral and octahedral coordination. The occupied octahedral sites were assumed to be present between sandwiches stacked in the 2H-TaS2 type of stacking of sandwiches: therefore no crystallographic shear has taken place during intercalation of these gaps. So far we assume the phase I a and 2a to be metastable phases which are found because of kinetic reasons. The EMF versus composition curve of the system Ag~TiS2 (where no shear transformations occur) measured at 300 K in a liquid electrolyte cell [ 3,17] is, apart from a shift due to the temperature dependence of the EMF, almost identical to that at 450 K from a solid electrolyte cell [ 16 ]. It is therefore reasonable to assume that the shear transformations in the case of ' ~ , x T a S 2 phases are responsible for the differences between the results of Scholz and Frindt [ 3 ] at room temperature and our measurements at about 450-500 K using samples prepared at high temperature. Left and Butz [ 161 measured the voltage of the cell under charge and decharge of 30 gA cm -2. Upon intercalation, their curve shows a large change in the EMF for 0.2 < x < 0.4; upon deintercalation the largest change in the voltage is for 0.3 < x < 0.5. Plateaus with constant voltage are hardly visible. Leffand Butz [18] also measured the ~S~Ta nuclear quadrupo~e frequency during intercalation and deimercaiation. These measurements allow to identify dis~,i~z~ inter° calated phases. Upon in*e~a~a~ion as well as dei~,~tercalation a pseudo-second stage phase is present for x < 0.4 with a maximum molfraction of silver of 2t340% (for intercalation and deintercalation respectively). Left and Butz ascribe this phase as due to a "diffuse intercalation front". The second stage phase, almost pure for 0.33 < x < 0 . 4 2 , is present over the

1119

whole compost:ion range indicating that a transition to the thermodynamically stable form is slow. The first stage phase, present for 0 . 4 2 < x < 0 . 6 7 upon intercalation and for 0.55
[ 3 ], Left and Butz [ 16 ] and ours agree in the region where the largest change in the EMF occurs i.e. for 0.2 < x < 0 . 4 , the homogeneity region for the second stage phase.

4. Discussion of the thermodynamical properties The thermodynamic aspects of intercalation cempounds of transition metal dichalcogenides have been discussed by several authors, mainly concerning the alkali metal intercalates: LixTiS2 [ 19,20], LixTaS2 [21 ], NaxTiS2 and NaxTaS2 [22 ]. The lithium intercalates LixTaS2 and LixTiS2 have a nearly linear ,~,,,-v,,~,-,,,,,~, va,,,~,,,, ,,, tl~e EMF in the range 0 < x < 1; the EMF decreases from about. 2.5 V for x = 0 to about 1.8 V for x = 1. One can find an expression for the x and T dependence of the EMF of a ceil M / M + conductor/ MxTX2 (M = Li,Na) in the case of single lrnase MxTX2 ( O < x < 1 ) from an expression for type free enthalp/ of intercalation ~.G (being the free enthalpy of formation of the intercalate from x moles M and one mole TX2 of the form: AG= - A x + (½)Bx 2+ R T X [xln x + ( 1 - x ) In ( 1 - x ) ] , in which A x represents the binding energy (A > 0); the term in x 2 originates from the repulsive intec-.~:ien betwee, intercalated atoms ( B > 0 ) . The charge in kermi energy due to filling of the conduction band is neglected or assumed to be linear in x2; the term ( ½)Bx 2 in that case also includes th~s variation. The term proportional to the temperature 7", representing the entropy, comes from mean fie~d theory: the intercalated atoms are assumed to be statistically distributed over the available sites. The EMF is giver by E=-AG/F=-(dAG/dx)/F =A/F-Bx/F-

( R T / F ) In x~ ( 1 - x ) ,

G.A. Wiegerset al./lntercalation compoundsAgxTaS.,

1120

in which F is the Faraday constant. A67 is the difference of the partial free enthalpy of the metal in the intercalate (G) and the free enthalpy of pure metal (G°). From the experimental EMF versus x curve one may find by integration the free enthalpy of intercalation X

AG= - F ~ E ( x ) d x . 0

The entropy of the intercalation AS, is found from

AS=Fi [dE(x)/dTldx, 0

F(dE/dT) being the difference between the partial entropy ~qof the metal in the intercalate and the entropy S ° of pure metal. The enthalpy of intercalation, AH, can be otained from AG= AH-TAS. Staging as observed in AgxTiS2 [2,15,17,18 ], Ag,NbS2 [ 14] and Ag~TaS2 is also present in sodium intercalated TiS2 [ 22 ] and TiSe2 [ 4 ]; a second stage phase was also reported for LixTaS2 [21 ]. In order to explain staging, besides in-plane interactions between intercalated atoms, also repulsive interplanar interactions were taken into account. For the graphite intercalates Safran [23] used a repulsion between atoms in different layers, the intralayer interaction between atoms was taken attractive. A generalization of Safran's approach was given by Millman and Kirzenow [ 24 ]; a model for the energy needed to expand a gap was given by Armand [25] and Dahn et al. [26]. For the stage-1 and -2 phases NaxTiS2 and NaxTiS2 with Na in trigonal-prismatic coordination, the trigonal-prismatic sites form a hc lattice, similar to the puckered hc lattice of tetrahedral sites in AgxTaS2. For the explanation of the EMF versus x curves Nagelberg and Worrell [22 ] took into account an inplane repulsive nearest neighbour interaction, the change in Fermi energy due to filling up the d-band with electrons from the alkali metal and an energy needed to expand the gap. The configurational part of the entropy was calculated with mean field theory, assuming ~hat sites at distances (1/3)a\/3 apart cannot be occupied simultaneously. For small values ofx (x< 0.2 ) the number of sites per cell is then ap-

proximately 2-3x since every occupied site excludes three neighboudng sites at ( 1/3)ax/~; above x=2,' 3 the number of available sites is one [ 22 ]. A major difference between the system Ag~TaS~ (and also of Ag~NbS2 and AgxTiS2) and the corresponding alkali metal intercalates, is the much lower value of the EMF and therefore also the free enthalpy of intercalation of the silver intercalates. One can understand this from the difference in the sum of the ionization and sublimation energy of silver compared to the alkali metals [27]. Another important difference between silver and alkali metal intercalates concerns the entropy; the temperature coefficient of the EMF is always positive for the silver intercalates and negative for the alkali metal intercalates (table 1 in ref. [ 20 ] ). The large positive value of A~qand AS for the silver intercalates cannot be explained by only the configurational entropy of the system; one has to include the entropy due to the lattice vibrations. For silver atoms coordinated by chalcogen atoms in layered compounds it has in many cases been found by X-ray diffraction that the vibration is anisotropic with large root-mean-square amplitudes in the layer-plane, e.g. AgCrS2 [28,29], stage-l and stage-2 AgxTiS2 [ 15,30], AgxNbS2 (isostructural with AgxTaS2 [ I l, 12] ). For AgCrS2, with silver atoms on a ha!f-filled hc lattice, Bruesch et al. [31 ] have found by means of far-infrared and inelastic neutron scattering that there are low-frequency acoustic and fiat optic modes. The silver ions, being strongly involved in these low=frequency modes, vi= brate parallel to the layers with large root-mean= square displacements. Long wavelength phonons were also observed in the intercalated AgxTaS2 [ 32 ]. A large r.m.s amplitude of vibration ( ( r 2) I/2) means a small value of the constant ce in the aarmonic potential V=ar 2 since ( r E) ~/2= k T / ( 2a ), which in turn corresponds with a small value of the vibration frequency. Whep one taken an Einstein model of vibrating atoms with frequency v' in silver metal and v in the intercalate, the change in the vibrational entropy is given by AS=3Rx In (v'/v), which is positive if v ' > v. in the expression for the free enthalpy of intercalation, a term proportional to Tx appears, which gives in the EMF a positive contribution independent of x For the system Ag,TiS: it was indeed observed that the partial entropy which is as high as 35-40 J/(mole K) does not change much

G.A. Wiegers et aL/Intercalation compounds Ag~TaS2

with composition [ 15,18 ]. For the system AgxTaS2 this was also observed. It means that the Einstein model for the description of the vibration of silver atoms in intercalates is not bad. For the system Ag~Ta~i2 ~ e free enthalpy of intercalation ziG, given in table 2, was obtained from the EMF curve of fig. 3. A rather crude approximation of the intercalation entropy was obtained assuming the temperature coefficient of the EMF to be 25 m V / K (A.q= 25 J tool- t K - t ) over the whole composition range. It is seen (table 2) that the entropy is an important factor in the stability. The same is the case for the system Ag~TiS2 [ 15,18 ]; the system "floats" on the entropy. Staging in the system AgxTaS2 can be understood in the same way as discussed for the alkali metal intercalates. The free enthalpy of intercalation of the first, second and third stage phase (AG(1), AG(2) and AG(3) respectively) is written as:

AG( l )= - a x + (½)Ox2+ (½)Cx +Jx/ (7+x)+RT[xlnx + (l-x)

In ( l - x ) l ,

for 0 < x < 1; A G ( 2 ) = - A x + B x Z + J x / ( y + 2 r ) + ( ~) R T X [ 2x In 2x + ( 1 - 2x) In ( 1 - 2x) ] , for 0 < x < 1/2; AG( 3 ) = - A x + ( 3 ) B x 2+ y x / (7+ 3x) + (-~)RT

X [3xln 3x+ ( 1 - 3 x ) In ( 1 - 3 x ) ] , for 0 < x < 1/3.

lance a from each other is accounted for by ( ½)Bx 2 for a first stage phase, ( I ) B (2x) 2= Bx 2 for a second stage phase and ( 1 ) ( ½) B (3x) 2= ( ~ ) Bx 2 for a third stage phase. The in-plane repulsion between atoms was taken the same for stage 1, 2 and 3 because of the similarity of the in-plane structure. In the case of the stage-1 phase one has to include a term (½)Cx 2 which accounts for the repulsion between silver atoms of neighbouring silver layers ( C < B because of the larger distance). The energy needed to expand the van der Waals gap is calculated in the rigid plate model of Dahn et al. [ 26 ]. The sandwiches TaS2 are assumed to be rigid plates connected by springs with spring constant K and equilibrium length Co, the thickness of a sandwich in 2H-TaS2. The elastic effects from silver are modeUed by springs with spring constant k and equilibrium length ct
The repulsion energy between silver atoms at dis-

Table 2 Thermodynamic properties of Ag,.TaS, from EMF measurements at 44 K.

0.10 0.22 0.30 0.35 0.56 0.67

AG (kJ/mol)

i'/.~5 (kJ/mol)

AJI (kJ/mol)

-2.5 -5.2 -6.7 -7.3 -9.0 -9.5

1.I 2.4 3.3 3.9 6.2 6.8

-1.4 -2.8 -3.4 -3.4 -2.8 -2.7

1121

x~

AG, J/tool

o -2000 ~ -4000

stage-3 stage-2 stage-1

-6000

..... : ~ : . . . . z.,:...

-8000

-I00001 0.o0

0.25

0.50

0.75

x Fig. 4. AG versus x curves of stage 1, 2 and 3 from a model, see texl.

G.A. Wiegerset al./IntercalationcompoundsAg.,.TaS2

1122 250

E~F, mV

2+3 ....................

.

200 150

';. •....

I+2

100

""\.... 50

'°N -

).00

'

'

'

'

I

'

0.50

0.25

0.75

X

Fi; ~. EMF of the electrochemical cell versus x from a model, S,~ t:; : .

AG versus x plots o f stage 1, 2 a n d 3 are caiculated using the expressions given a b o v e with: A = 4 0 k J, B=36 kJ, C= 16 kJ, J= 10 kJ, y = 0 . 1 (fig. 4). T h e c o m p o s i t i o n s o f coexisting p h a s e s can be f o u n d graphically f r o m the c o m m o n t a n g e n t o f the G curves. More accurate values o f the c o m p o s i t i o n s o f coexisting phases can be f o u n d by calculating the g r a n d potential O=AG-xAG for each p h a s e a n d drawing plots o f ~ against AG; coexisting phases have the same value o f O a n d o f A(~. The p a r a m e t e r s were found by trial a n d error; with the p a r a m e t e r s given above the E M F versus x plot (fig. 5), c o r r e s p o n d i n g with the AG versus x plots o f fig. 4, shows the essential features o f the e x p e r i m e n t a l E M F versus x curve o f fig. 3. There are single phase regions separated by two phase regions, stage 1 and stage 2, stage 2 and stage 3, a n d a t w o - p h a s e region o f stage 3 a n d a 2 H - T a S z phase with a very small a m o u n t o f silver. The phase limit~ a n d also the E M F at the p l a t e a u s do not correspond q u a n t i t a t i v e l y with the e x p e r i m e n t a l values; this is not unexpected since the m o d e l is rather crude. The h o m o g e n e i t y region o f the a s s u m e d stage-3 region is r a t h e r small and with slightly different parameters (especially J and y are i m p o r t a n t ) the stage3 phase does n o t occur because a two-phase m i x t u r e of stage 2 a n d 2 H - T a S 2 is m o r e stable. It might ex-

Re_~erences [l ] J.M. van de Berg, Thesis (University of Leiden, The Netherlands, 1964). [2] G.A. Scholz and R.F. Frindt, Mater. Res. Bull. 15 (1980) 1703.

[3] G.A. Scholz and R.F. Frindt, J. Electrochem. Soc. 131 (1981) 1763. [4] H.J.M. Bouwmeester, E.J.D. Dekker, K.D. Brorsema, R.J. Haange and G.A. Wiegers, Rev. Chim. Miner. 19 (1982) 33. [ 5 ] G.A. Scholz, R.F. Frindt ~nd A.E. Cu~on, Phys. Status Solidi (a)71 (1982) 531; (a)72 [1982) 375. [6] G.J. Tatlock, E.A. Marseglia and R.H. Friend, Int. Phys. Conf. Ser. no. 68 (Guildford, 1983) 55. [ 7 ] G.S. Boebinger, N.I.F. W~kefileld, E.A. Marseglia and R.H. Friend, Physica 117B/118B (,1983) 608. [ 8 ] A. Diedoring, R.J. Haange, G.A. Wiegers, J. Mahy, J. van Landuyt and S. Amelinckx, in: 3rd European Conf. on Solid State Chemistry, (1986) p. 263. [9] H. Sate and R. Kikuchi, J. Chem..D,~ys. 55 ( 1971 ) 677. [ 10] K.R. Subbaswamy and G.D. Mahan, Phys. Rev. Letters 37 (1976) 642. [ 11 ] J. Mahy, G.A. Wiegers, E v o n Bolhuis, A. Diedering and R.J. Haange, Phys. Status Solidi Ca) ( 1988 ), to be published. [ 12 ] H.J.M. Bouwmeester, Theses (Groningen, 1988). [ 13] W. Weppner and R.A. Huggings, J. Electrochem. Soc. 124 (1977) 156. [ 14] H.LM. Bouwmeester, Solid StaLe ion:~ 16 (19°5) 163. [ 15 ] A.G. Gerards, H. Roede, R.J. Haange, B.A. Boukamp and G.A. Wiegers, Synth. Metals I:~ (1984/85) 51. [ 16 ] A. Left and T. Butz, in: Reactivity of solids, eds. P. Barret and L.C. Dufour (Elsevier, Amsterdam, 1986 ). [ 17 ] J. Schranke and R. SchSllhorn, Solid State Ionics 23 (1987) 197. [ 18 ] A. Honders, E.W.A. Young, A.H. Heerens, J.H. de Wit and G.H.J. Broers, Solid State Ionics 9/l 0 (1983) 375. [19]A.J. Berlinsky, W.G. Unruh, W.R. McKinnon and R.R. Haering, Solid State Commun. 31 (1979) 135. [20] A. Honders, J.M. der Kinderen, A.H. van Heeren, J.H.W. de Wit and G.H. Broers, Solid State r,onics 14 (1984) 205. [21 ] JR. Dahn and W.R. McKinnon, J. Phys. C 17 (1984) 4231; J. Electrochem. Soc. 131 (1984) 1823. [22] A.S. Nagelberg and W.L Won'ell, J. Solid State Chem. 29 (1979) 345; 38 ( 1981 ) 321. [23] S.A. Safran, Phys. Rev. Letters 44 (1980) 937. [24] S.E. Millman and G. Kirzenow, Phys. Rev. B 28 ',19~) 3482. [ 25 ] M. Armand, Thesis (University of Grenoble, 1978). [26] J.R. Dahn, D.C. Dahn and R.R. Haering, Solid State Commun. 42 (1982) 179. [ 27] G.A. Wiegers, H.J.M. Bouwmeester and A.G. Gerards, Solid State lonics I0 (1985) 155. [28] A.G. Gerards, B.A. Boukamp and G.A. Wiegers, Solid State :vnics 9/10 (1983) 1lCJ3. [29] A.G. Gerards, A. Diedering, R.J. Haange a'~d G.A. Wiegers, to be puL!ished. [30] G.A. Wiegers, K.D. Eronsema. S. van Smaa;: ", R.I. Haange, LE. Zondag and .I.L de Boer, J. Solid State Chem. 67 ( 1987 ) 7. [31]P. Bruesch, T. Hibma and W. Buhrer, Phys. Rev. B 27 ( 1983 ) 5052. [32] W.G. McMullaa and J.C. Irwin, Can. J. Phys. 62 (1984) 789.