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The ionic conductivity variation in rapidly quenched lithium-containing glasses
Solid State lonies 2 (1981) 163-170 North-Holland Publishing Company
THE IONIC CONDUCTIVITY VARIATION
IN RAPIDLY QUENCHED LITHIUM-CONTAINING GLASSES K. NASSAU, R.J. CAVA Bell Laboratories, Murray Hill, NJ 07974, USA and
A.M. GLASS Bell Laboratories, Holmdel, NJ 07733, USA Received 14 November 1980 The lithium ion conductivity of a wide variety of rapidly quenched glasses is studied both as a function of lithium ion content and with various additives which are likely to affect the microstructure of the glass. In each of the glass systems studied a maximum ionic conductivity of R,10-a (f~ cm)-1 is observed at 500 K, but this value is reached for different Li ion concentrations in different systems. Experiments with additions to the glass composition suggest that the availability of vacant interstitial sites in glasses of this type is not a limitation to fast ion conduction.
1. Introduction
Several factors have been shown to be of importance in the development of interstitial-type superionic conduction in crystalline materials. These fall into two categories: those affecting the ionic mobility, and those affecting the proportion of the conducting ionic species which are mobile. The ionic mobilities are dependent on chemical bonding effects, size of the mobile ion relative to its diffusion pathway and ion-ion interaction. The proportion of ions available for conduction is determined by stoichiometry, and is most importantly dependent on the number of vacant interstitial sites of roughly equivalent energy per mobile ion. We have previously reported on fast ion conduction in rapidly quenched oxide glasses [1-5] and have attempted to change the mobile ion concentration with respect to available interstitial sites based on arguments that are sound in classical crystal chemistry, but cannot be easily backed up by structural arguments for non-crystalline solids. For instance we have studied the effect of h'thium ion concentration on the ionic conductivity of Li20 : A1303, Li20 : Ga203 and Li20 : Bi203 glasses [4]. An exponential
behavior of the ionic conductivity with Li ion content was observed. On the other hand, similar experiments with Li20 : WO 3 and Li20 : MoO 3 showed relatively little variation of conductivity with Li ion concentration [5]. Furthermore, the conductivity of the Li20 : Nb205 system over a limited composition range reaches a maximum value at ~50 cation percent lithium [2]. This multiplicity of behavior indicates that no single explanation can account for all the observed behavior. In this report we examine the ionic conductivity variation in lithium-containing glasses in more detail. The results on several rapidly quenched glasses are compared with the results obtained previously with glass-forming systems. More extensive data on the Li20 : Nb205 system, and limited results on quenching the L i 2 0 - T a 2 0 5 system are presented. Attempts to increase the ionic conductivity with glass formers or by chemical additions which would increase the "vacancy" or "non-bridging" oxygen ion concentration according to classical argument in crystalline or glass-forming systems are reported. Because of the absence of any structural information with which to prodiet the behavior of these rapidly quenched glasses, such experiments can be useful to give insight into cer. tain structural characteristics of the glass.
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I(. Nassau et al. / Ionic conductivity variation in rapidly quenched lithium glasses
164 2. Experimental
Details on roller quenching with a quench rate of ~-107 degrees per second and evaluation o f the presence of glass by examination of the resulting flakes by differential thermal analysis (DTA) and powder X-ray diffraction, have been described in detail elsewhere [ 1]. The lithium concentration is given as the cation fraction o f Li = [Li]/([Li] + ~ [ M ] ) where M represents all other cations present. The cation fraction defined in this way is to be preferred to the mole fraction o f lithium salt since the density of mobile lithium ions is the important parameter for ionic conduction. Ionic conductivity measurements were all made using gold electrodes, which are blocking to lithium ions, evaporated onto the glass flakes. The sample resistance was obtained on selected samples b y measuring the complex ac impedance over the frequency range from 100 Hz to 1 MHz. The electrode polarization characteristic was extrapolated to zero imaginary part in the standard complex impedance plot. This usually occurred at frequencies near 500 kHz for the sample geometries and resistances encountered. Other
samples were studied using automatic 1 MHz and 1 kHz impedance bridges for rapid measurement as a function of temperature.
3. Results 3.1. L i 2 0 - N b 2 0 5 and L i 2 0 - T a 2 0 5 The experimental data for the 16 niobate and 4 tungstate compositions quenched are presented in tables 1 and 2. Glass in variable amounts was obtained from Li = 0 . 3 0 - 0 . 7 5 in the niobate and at 0.50 and 0.60 in the tungstate systems with Li20. In fig. 1 is shown the phase diagram of the L i 2 0 N b 2 0 5 system based on various studies as summarized by Levin and McMurdie [6]. Superimposed on this figure are the initial crystallization temperatures (circles) o f the glasses o f table 1; arrows indicate compositions where glass could not be obtained. Throughout the whole glass-forming region, LiNbO 3 (JCPDS 9.186) [6,7] is the first phase to crystallize, and the minimum at the LiNbO 3 composition can accordingly be readily understood. Since a diffusion process is,
Table 1 Quenched specimens in the Li20-Nb20 s system Li content a)
Compound (if known)
Number of quenches b) and temperature CC)
Occurrence of glass c) Crystallization and temperature (°C) exotherm d) (oC)
0.067 0.10 0.20 0.25 0.30 0.385 0.435 0.465 0.488 0.50 0.512 0.556 0.60 0.70 0.75 0.80
LiNb2aOTt LiNb30 s LiNbO3 LisNbO 4 -
1; 1400 2; 1600 1; 1500 2; 1750-1850 1; 1475 1; 1330 1; 1350 3; 1300-1550 6; 1475-1600 15; 1325-1700 2; 1425-1525 5; 1250-1650 2; 1500-1590 2; 1550 2; 1550-1700 2; 1550-1700
trace; 1475 pure, 1330 some, 1350 pure, 1550 pure, 1550 pure, 1550 much, 1525 some, 1500 trace, 1500 trace, 1550 trace, 1700 -
a) All compositions were sintered at 1000°C. b) The number of separate melts which were quenched at each composition. c) All glasses were colorless; amount of glass in best preparation is listed with the quenching temperature. d) For each instance the first phase to crystallize was LiNbO3.
~600 520 510 470 460 460 460 630 ~650 ~700 ~710 -
K. Nassau et aL / Ionic conductivity variation in rapitffy quenched lithium glasses
165
of necessity,involved,leastdiffusionis required when
1200 w f,. p_ 1000
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,-LINbO 3 /LiNb308
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800
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0 NblO 5
•
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I 0.4
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I
t' 0.8
0.6
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t.0 L i 20
LI (CATION FRACTION)
Fig. 1. Phase diagram of the Nb2Os-Li20 system with circles showing the crystallization temperatures of quenched glasses. Arrows indicate compositions where glasses could
not be obtained.
.4R
t0-3
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/'.
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I 0.2
I 0.4
I 0.6
I 0.8
1.0
Li20
ki (CATION FRACTION)
Fig. 2. Ionic conductivity variation with Li contant in quencheal glassesat 500 K for LiTaOa and the Nb2Os-Li20 system.
the composition o f the glass is the same as that of the crystallizing phase. The further away from LiNbO3, the more excess Li20 or Nb205 is present, thus implying both a longer diffusion path, and possibly a less favorable environment for diffusion. The ionic conductivity of the glasses is highest at Li = 0.513, as shown in fig. 2, and falls off rapidly in either direction. The conductivity versus 1/T curves of some of these niobate compositions, as well as of LiTaO 3 also included in fig. 2 have been published previously [2]. Quenched 6LiNbO 3 glass made from atomic weight 6 isotope of Li had the same ionic conductivity characteristics as the ordinary LiNbO3, within the experimental error. Sargeant and Roy quenched both Ta205 and Nb205 to obtain metastable high-temperature phases, but no glass was observed [8]. In general, the quenching and glass characteristics in the L i 2 0 - T a 2 0 s system of table 2 appear to be similar to those of the L i 2 0 - N b 2 0 5 system of table 1. The L i 2 0 - T a 2 0 5 phase diagram [9] shows the existence of three compounds: LiTa6016 , Li3TaO4, and LiTaO3, the latter with some solid solution similar t o that of LiNbO 3 in fig. I. Tables 1 and 2 indicate that essentially pure glass could be obtained in these systems close to the Li = 0.50 concentration. Accordingly, the rapid fall-off in both directions in fig. 2 may be partially explained by a reduction in the amount of glass present with a concentration departure from Li = 0.50 in either direction. This may not be the full explanation, since a quantity of glass giving a conductivity down 2~ orders of magnitude at Li = 0.75 would probably not be detectable by DTA. Since glass is likely to be present as a continuous phase surrounding crystallites, a large reduction in conductivity resulting from discontinuities in the electrical conductivity path would not be expected. It may be noted that the ionic conductivity of single crystal LiNbO 3 extrapolated to room temperature is ~1020 times smaller than that of quenched vitreous LiNbO 3. In the mixed LiNbO3-LiTaO 3 system glass was obtained throughout the whole system and the crystallization temperature varied monotonically with composition: from 100% LiNbO 3 to 100% LiTaO 3 there was very little observable change on the ionic conductivity. All of this indicates a smooth "isomor-
166
If. Nassau et aL / Ionic conductivity variation in rapidly quenched lithium glasses
Table 2 Quenched specimens in the Li20-Ta2Os system Li content a)
Number of quenches and temperature (°C)
Occurrence of glass b) Crystallization and temperature (*C) exotherm c) (0C)
0.33(LiTa30 a 0.50(LiTaOs) 0.60 0.80
1; 1800 7; 1685-1890 3; 1700-1900 3; 1700-1850
much, 1850 much, 1800 -
570-630 620, 650 -
a) All compositions were sintered at IO00°C. b) All glasses were colorless; amount of glass in best preparation is listed with the quenching temperature. c) For each instance the first phase to crystallize was LiTaOa. phous" substitution of Nb and Ta for each other in the vitreous phase. This is definitely not the case in the mixed LiNbO 3 : KNbO 3 : NaNbO 3 system [10] where a marked reduction of the conductivity is observed when two alkalis are present in the glass composition as shown in fig. 3. The reduction in conductivity on the mixing o f alkali ions is a phenomenon well known in the network-forming glasses such as silicates, as the "mixed alkali effect" [11]. The activation energy shows an analogous behavior [ 10]. 3.2. Variation o f ionic conductivity with lithium con ten t
In the most general form, the lithium ion conductivity o can be expressed as o = ~
i
"fui,
I0-~
Ss "q
"7
~
i
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r
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,
,~ ~o-~
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i
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I I ~s
0'2 d , NoNbO 3
LiNbO 3
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/
0 .'6 0 .'8
KNbO 3
NoNbO s
Fig. 3. The mixed alkali effect in LiNbO s : K N b O 3 : NaNbO a q u e n c h e d glasses.
where the sum is over all types of Li ion sites i, n i is the density of ions of charge e at sites i, and #i is a mobility factor which contains the details o f the structural environment of site i. In a glass, some Li ions may be immobile with/a i = 0, while others may move freely from one site to another. Thus to maximize the ionic conductivity involves maximizing the density of mobile ions and their mobility. In general, increasing the Li ion concentration o f a glass will also change the mobility, so that the net result may not be an increase of a. In this section experiments are described which were designed to study the effects on o of varying the Li ion content in various glass systems, while attempts to vary the microstructure by chemical addition to the glass in such a way as to increase the mobility are described in the following section. The ionic conductivities at 500 K of the quenched glass systems of Li20 with Nb205, Ta205, A1203, Bi203, Ga203, WO 3 and MoO 3 are shown in fig. 4. Also included are the data from the quenched mixed sulfate glasses in the system Li2SO 4 with SrMg(SO4) 2 and La2(SO4) 3 [12,13], and from some mixed lithium glasses containing B20 3 and other ingredients as published by Smedley and AngeU [14]. The mole fractions quoted in this last work have been recalculated as cation fractions. Three characteristics can immediately be observed in fig. 4. First, there is similar behavior among equivalent valence cations which might reasonably be expected to substitute for each other, even when there is a large variation of ionic radius as that in going from the 0.53 A o f A13+ to the 1.02 A of Bi 3+. Second, there are important differences in the shapes of the curves; there are also differences in the location of
K. Nassau et aL / Ionic conductivity variaffon in rapidly quenched lithium glasses -2
.$ "7
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I
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IN "S
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0.2
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,
0.8
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Fig. 4. Ionic conductivity variation with Li content at 500 K for Li-containingoxide and sulfate glasses. glass-forming regions. Third and most surprising, the best ionic conductivities in each system are all about the same, ~10 - 3 ( ~ cm) -1 even though they occur at different Li concentrations. It is clear from these data that no simple explanation is likely to account for all the observations. The exponential increase of conductivity with lithium ion content in the case of the Li20 : A1203, Li20 : Ga203 and Li20 : Bi203 systems can be acounted for in terms of a constant ion mobility and an exponential increase of the mobile carrier concentration, even though the total lithium content only increases linearly. This explanation has been used to account for the increase of o with Na content in the SiO 2 : Na20 system [15] where the Na ions are assumed to be distributed between two types of sites: those which are associated with non-bridging oxygen ions and those which are dissociated (the weak electrolyte model). However, in the absence of any structural information in these glasses an equally plausible explanation for the exponential increase of o can be given in terms of a large number of alkali ion sites having a gaussian distribution of activation energies [4]. Within this picture it is not possible to distinguish mobile and immobile lithium ions. An approximately exponential increase of conduc-
167
tivity with lithium contant is also observed with the Li2SO 4 : La2(SO4) 3 system [12], although the exponent is only about one third of the value for the oxides just discussed. In terms of the random site mod. el this implies a narrower distribution of site energies, whereas in terms of the weak electrolyte model this would imply a smaller energy difference between associated and dissociated lithium ions. In either case this might be expected for sulfates since the negative charge on the (SO4) 2 - ion is distributed over a wider volume and would have a smaller binding energy to the lithium ions. The conductivity of the Li20 : w e 3 and Li20 : MoO 3 systems is relatively constant as the lithium ion concentration is increased. In this case, it would appear that Li ion mobili W is relatively unaffected by the composition - in agreement with studies of Li diffusion in electrochromic w e 3 films. In the Li20 : Nb205 and Li20 : Ta205 systems a pronounced maximum of the conductivity is observed for a cation fraction near 0.5. As discussed earlier, it is not likely that this is due to the formation of crystalline material within a glass since it is known that a small amount of crystallized material has little effect on the conductivity. The most reasonable explanation of the maximum is that the lithium ion mobility reaches a maximum value for 50 cation percent lithium and a further increase results in a decreased mobility. This can arise from the formation of negatively charged centers such as non-bridging oxygen ions which act as alkali ion traps in the structure. We have already pointed out tliat if the structure consists of a random array of corner linked NbO 6 octahedra, then the concentration of non-bridging oxygen ions is minimum (zero) for x = 0.5. A similar argument would predict maxima in Li20 : MgO and Li20 : T i e 2 glass systems at compositions of Li4MgO 3 and LiTiO 3 if the Mg and Ti ions are octahedrally coordinated. We have, however, been unable to prepare glasses of these compounds to verify this model. In both the Li2SO4 : SrMg(SO4) 2 glass system [13] and the mixtures reported by Smedley and AngeU [14] a decrease of the ionic conductiviW is observed with increasing Li ion concentration for high concentration values. In these cases the systems are too complex to speculate about the origin of the behavior.
168
K. Nassau et aL / Ionic conductivity variation in rapidly quenched lithium glasses
Table 3 Quenched alkali niobate and tantalate specimens with the addition of glass-formers or "vacancy" inducers Composition a)
x
LiNbOs LiNb t-xOs72.sx*xSi02 LiNbt_xOs_2.sx.xGe02 LiNbt_2xO3_sx*xB203 (1-x)LiNb0a.xLAS
0.2 0.2 0.1 0.1 0.2 0.1 0.02 0.1 0.2 0.05 0.1 0.03 0.1 0.2 0.1 0.2
(1-x)LiNbO 3 .xSLS (Lil_ 2xMgxe)x)Nb03 (Li I _ 3xAlx¢2x)Nb03
(Lil-x¢x)Nbt-xWx03 Li(Nb 1-2xWxCx) (03-2x¢2x) KNbO3 (K t - 2x Lax~2x)Nb0 3 LiTaO 3 Li(Tat_ ~xWxCx)(O3_ 2x~2x)
-
0.1 0.1 0.2
Number of quenches and temperature (°C) 15; 1325-1700 1; 1800 1; 1750 1; 1750 1; 1550 1; 1300 1; 1560 6; 1500-1800 5; 1350-1700 5; 1225-1500 1; 1550 1; 1550 4; 1475-1750 1; 1350 1; 1300 1, 1400 1; 1375 3; 1175-1310 1; 1450 7; 1685-1950 1; 1600 1; 1525
Occurrence of glass b) and exother~ (*C) pure, 460 pure, 550 c), 610 pure, 530 e), 580 much, 510 pure, 520 c) 560 pure, 620 pure, 520 c),~60L much, 460 pure, 470 some, 500 pure, 540 pure, 550 pure, 460 pure, 460 pure, 450 much, 490 some, 480 pure, 380 e), 460 pure, 620 some, 570 some, 550 some, 610
a) LAS is the lithia alhmina silicate 35 : 5 : 60; SLS is the soda lime silicate 20 : 11 : 60; ~ indicates a "vacancy". b) Amount of glass in best preparation is listed; only strong, initial crystallization exotherm is given. c) Glass transition endotherm.
3.3. The variation o f ionic conductivity with additives In the crystalline state alkali niobates and tantalates are most frequently characterized b y the corner linking o f NbO 6 octahedra in an infinite threedimensional array: for instance the perovskite (KNbO3, NaNI~O3) , tungsten bronze (Ba4Na2Nbl0030 , K6Li4Nbl0 O30), or LiNbO 3 structures are all characterized in this way. The alkali ions occupy the interstices between the NbO 6 octahedra. In the case of the perovskites there are no vacant interstitial cation sites in the structure available for the m o t i o n of alkali ions. In LLNbO 3 and the tungsten bronzes there are vacant cation sites, but the structural barriers between these sites prevent fast ion motion. The basic building blocks o f the glass are again likely to be NbO 6 octahedra, corner shared, but because of the random structure of the glass there exist paths which allow m o t i o n of alkali ions through-
out the structure which do n o t exist in the crystalline state. The question we attempted to answer was whether this ionic mobility for the glass could be increased b y introducing additional vacant cation sites b y doping with cations o f different valence. This procedure is of course well known to increase the ionic mobility in many crystalline materials. We have already shown that the addition o f MgO and A1203 to LiNbO 3 glass results in decreased conductivity, while the addition o f WO 3 produced little change o f the conductivity [3]. However, within the above model for the glass structure these additions would result in the creation o f negatively charged non-bridging oxygen ions in the glass which act as traps for the Li ions. To avoid this, KNbO 3 glass was doped with L a N b 3 0 9 , which would retain the corner linked NbO 6 framework without introduCing nonbridging oxygen ions while increasing t h e d e n s i t y o f
K. Nassau et aL / Ionic conductivity variation in rapidly quenched lithium glasses
vacant interstitial cation sites, i.e. the composition employed is most illustratively represented by
(Lil.8550.15)(Wl.85 Nb0.15)O6.85,
KO.7 Lao. 1 ~0.2 NbO 3 ,
(Lil.7~b0.15 Mg0.15)W207,
where ~ indicates a vacancy. This and other such compositions quenched are listed in table 3. The conductivity of this glass at 500 K was measured to be 2 × 10 -7 (I2 cm) -1 - about two orders of magnitude lower than previously reported for a iaure KNbO 3 glass. If the structural picture is correct, this resuli imp]i6:~'tl~at the number of vacant alkali ion sites in thepure KNbO 3 glass is already sufficient for fast ion conduction, and the reduction in conductivity is due to the reduction of the total density of alkali ions available for transport. It was shown in the preceding section how the conductivity can be an extremely non-linear function of alkali ion concentration. Of course, there may be other effects present, such as blocking of the tunnels available to K transport by the large La 3+ ions, or changes in the basic structural framework. A sirnflar set of substitutions in Li2W20 7 glass was performed since in this system the ionic conductivity is relatively insensitive to the alkali content. The glass structure is probably built upon comer linke d WO4 tetrahedra with Li ions at the interstices. We attempted to increase the density of vacant interstitial sites via the substitution of Mg2+ or A13+ for Li+, or W6+ for Nb 5+ ions. The compositions employed are shown in table 4 and are
169
(Li 1.7~0.2A10.1)W207 • Remeasurement of the Li2W207 glass indicated significant electrode polarization at 1 kHz, where previous measurements were made, and an actual conductivity at 500 K of ~-3.5 X 10- 3 ([2 cm) -1 , one order of magnitude greater than that reported from the 1 kHz measurements. This conductivity compares favorably with the best Li ion conductivities reported for glasses and also that for Li3N and Li~-alumina. All three additions to Li2W207 resulted in a decrease in the conductivity. For both the A1 and the Mg additions, a at 500 K was ~1.0 X 10 - 3 ([2 cm) -1 , and for the Nb/205 additions o at 500 K was slightly high. er, ~1.5 X 10 - 3 (~2 cm) - 1 . These decreases are small compared to that which resulted from the addition of Lal/3NbO 3 to KnbO 3. The effect on the ionic conductivity of network forming additives such as SiO2, GeO 2 and B203 to the Li20 : Nb205 and Li20 : WO 3 glass systems were studied for several compositions listed in tables 3 and 4. These additions of coruse facilitate glass formation and could have practical value if the mixed glasses had increased conductivity. In all cases, the addition of up to 20% of the glass formers produced no significant change of the ionic conductivity.
Table 4 Quenched lithium tungstate compositions with the addition of glass formers and "vacancy inducers" Composition a)
x
Number of quenches Occurrenceof glass b) and temperature CC) and exotherm (°C)
Li2W207 (1-x)Li2W2OT.xSiO 2 (1-x)Li2W2OT.xGeO2 (1-2x)Li2W2OT-2xB203 Li2W2_xO7_ 3x.xSiO2 Li2W2_xOT_ 3x.xGeO2 Li2W2_xO~_6x.xB203 (Li2_2xMgxc~x)W~07 (Li2_ 3xAlx¢2x)W207 Li2Wz_2xNb~x(OT_xCx)
0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.15 0.1 0.15
2; 925-975 1; 1350 1; 1290 1; 1300 1; 1350 1; 1250 1; 1275 1; 1210 1; 1170 1; 1180
pure, 340 c), 370 pure, 360 e), 390 pure, 420 c), 450 pure, 380 c), 410 pure, 350 e), 370 pure, 390 c), 440 pure, 370 c), 410 pure, 380 c), 400 pure, 380 e), 400 pure, 370 c), 400
a) LAS is the lithia alumina silicate 35 : 5 : 60; SLS is the soda lime silicate 20 : 11 : 60; ~ indicates a "vacancy". b) Amount of glass in best preparation is listed; only strong, initial crystallization exotherm is given. c) Glass transition endotherm.
170
IC Nassau et al. / Ionic conductivity variation in rapidly quenched lithium glasses
4. Conclusions
Acknowledgement
The most surprising result of this survey is that in each of the glass systems studied the same maximum value of the ionic conductivity, ~10 - 3 (~2 cm) -1 at 500 K was observed. However, in different glass systems the maximum occurs for different values of the lithium ion concentration indicating that no single explanation is appropriate for all glasses. Attempts to increase the ionic conductivity either by the addition of glass formers or by the addition of ions of different valence from those of the host glass were unsuccessful. These experiments indicate that the availability of vacant interstitial sites in glasses of the type studied here is not a limitation to the development of fast ion conduction. Of course, whether the atomic substitutions we have attempted resulted in vacant inter. stitial sites at all is open to question. It is dependent in part on the degree of short-range order and in part on the distance over which the concept of local charge neutrality can be expected to apply. The best way to view these glasses at present is as quenched liquids in which the liquid structure is largely preserved and, unfortunately, very little is known about the structure of these types of high melting point liquids. Concepts such as "non-bridging oxygen ions" or "vacancies" might not be appropriate in these types of solids.
We want to thank M. Grasso for assistance with many of the glass preparations and D.H. Olson for many of the conductivity measurements.
References [ 1] K. Nassau, C.A. Wang and M. Grasso, J. Am. Ceram. Soc. 62 (1979) 74, 503. [2] A.M. Glass, K. Nassau and T.J. Negran, J. AppL Phys. 49 (1978) 4808. [3] K. Nassau, M. Grasso and A.M. Glass, J. Non-Cryst. Solids 34 (1979) 425. [4] A.M. Glass and IL Nassau, J. Appl. Phys. 51 (1980) 3756. [5] IC Nassau, A.M. Glass, M. Grasso and D.H. Olson, J. Am. Ceram. Soc.127 (1980) 2743. [6] E.M. Levin and H.F. McMurdie, Phase Diagramsfor Ceramicists, 1975 Supplement (Am. Ceram. Soc., Columbus, 1975) p. 86. [7] JCPDS Powder Diffraction File, International Center for Diffraction Data, Swarthmore, PA (1977). [8] P.I. Sargeant and R. Roy. J. Am. Ceram. Soc. 50 (1967) 500. [9] F.I. Kozhaev, S.A. Shvartsman, S.I. Sadova and V.A. Sokholov, Opt. Mekh., Prom-st 45 (1978) 35. [10] A.M. Glass, K. Nassau and D.H. Olsen, in: Fast ions transport in solids, eds. P. Vashishta, J.N. Mundy and G.IC Shenoy (North-Holland, Amsterdam, 1979) p. 707. [11] J.O. Isard, J. Soc. Glass 43 (1959) 113. [12] K. Nassau and A.M. Glass, J. Non-Cryst. Solids (1981), to be published. [13] K. Nassau, A.M. Glass and M. Grasso, to be published. [14] S.I. Smedley and C.A. Angell, Solid State Commun. 27 (1978) 21. [15] D. Ravaine and J.L. Souquet, Phys. Chem. Glasses 18 (1977) 27.
THE IONIC CONDUCTIVITY VARIATION
IN RAPIDLY QUENCHED LITHIUM-CONTAINING GLASSES K. NASSAU, R.J. CAVA Bell Laboratories, Murray Hill, NJ 07974, USA and
A.M. GLASS Bell Laboratories, Holmdel, NJ 07733, USA Received 14 November 1980 The lithium ion conductivity of a wide variety of rapidly quenched glasses is studied both as a function of lithium ion content and with various additives which are likely to affect the microstructure of the glass. In each of the glass systems studied a maximum ionic conductivity of R,10-a (f~ cm)-1 is observed at 500 K, but this value is reached for different Li ion concentrations in different systems. Experiments with additions to the glass composition suggest that the availability of vacant interstitial sites in glasses of this type is not a limitation to fast ion conduction.
1. Introduction
Several factors have been shown to be of importance in the development of interstitial-type superionic conduction in crystalline materials. These fall into two categories: those affecting the ionic mobility, and those affecting the proportion of the conducting ionic species which are mobile. The ionic mobilities are dependent on chemical bonding effects, size of the mobile ion relative to its diffusion pathway and ion-ion interaction. The proportion of ions available for conduction is determined by stoichiometry, and is most importantly dependent on the number of vacant interstitial sites of roughly equivalent energy per mobile ion. We have previously reported on fast ion conduction in rapidly quenched oxide glasses [1-5] and have attempted to change the mobile ion concentration with respect to available interstitial sites based on arguments that are sound in classical crystal chemistry, but cannot be easily backed up by structural arguments for non-crystalline solids. For instance we have studied the effect of h'thium ion concentration on the ionic conductivity of Li20 : A1303, Li20 : Ga203 and Li20 : Bi203 glasses [4]. An exponential
behavior of the ionic conductivity with Li ion content was observed. On the other hand, similar experiments with Li20 : WO 3 and Li20 : MoO 3 showed relatively little variation of conductivity with Li ion concentration [5]. Furthermore, the conductivity of the Li20 : Nb205 system over a limited composition range reaches a maximum value at ~50 cation percent lithium [2]. This multiplicity of behavior indicates that no single explanation can account for all the observed behavior. In this report we examine the ionic conductivity variation in lithium-containing glasses in more detail. The results on several rapidly quenched glasses are compared with the results obtained previously with glass-forming systems. More extensive data on the Li20 : Nb205 system, and limited results on quenching the L i 2 0 - T a 2 0 5 system are presented. Attempts to increase the ionic conductivity with glass formers or by chemical additions which would increase the "vacancy" or "non-bridging" oxygen ion concentration according to classical argument in crystalline or glass-forming systems are reported. Because of the absence of any structural information with which to prodiet the behavior of these rapidly quenched glasses, such experiments can be useful to give insight into cer. tain structural characteristics of the glass.
0 167-2738/81/0000-0000/$ 02.50 © North-Holland Publishing Company
I(. Nassau et al. / Ionic conductivity variation in rapidly quenched lithium glasses
164 2. Experimental
Details on roller quenching with a quench rate of ~-107 degrees per second and evaluation o f the presence of glass by examination of the resulting flakes by differential thermal analysis (DTA) and powder X-ray diffraction, have been described in detail elsewhere [ 1]. The lithium concentration is given as the cation fraction o f Li = [Li]/([Li] + ~ [ M ] ) where M represents all other cations present. The cation fraction defined in this way is to be preferred to the mole fraction o f lithium salt since the density of mobile lithium ions is the important parameter for ionic conduction. Ionic conductivity measurements were all made using gold electrodes, which are blocking to lithium ions, evaporated onto the glass flakes. The sample resistance was obtained on selected samples b y measuring the complex ac impedance over the frequency range from 100 Hz to 1 MHz. The electrode polarization characteristic was extrapolated to zero imaginary part in the standard complex impedance plot. This usually occurred at frequencies near 500 kHz for the sample geometries and resistances encountered. Other
samples were studied using automatic 1 MHz and 1 kHz impedance bridges for rapid measurement as a function of temperature.
3. Results 3.1. L i 2 0 - N b 2 0 5 and L i 2 0 - T a 2 0 5 The experimental data for the 16 niobate and 4 tungstate compositions quenched are presented in tables 1 and 2. Glass in variable amounts was obtained from Li = 0 . 3 0 - 0 . 7 5 in the niobate and at 0.50 and 0.60 in the tungstate systems with Li20. In fig. 1 is shown the phase diagram of the L i 2 0 N b 2 0 5 system based on various studies as summarized by Levin and McMurdie [6]. Superimposed on this figure are the initial crystallization temperatures (circles) o f the glasses o f table 1; arrows indicate compositions where glass could not be obtained. Throughout the whole glass-forming region, LiNbO 3 (JCPDS 9.186) [6,7] is the first phase to crystallize, and the minimum at the LiNbO 3 composition can accordingly be readily understood. Since a diffusion process is,
Table 1 Quenched specimens in the Li20-Nb20 s system Li content a)
Compound (if known)
Number of quenches b) and temperature CC)
Occurrence of glass c) Crystallization and temperature (°C) exotherm d) (oC)
0.067 0.10 0.20 0.25 0.30 0.385 0.435 0.465 0.488 0.50 0.512 0.556 0.60 0.70 0.75 0.80
LiNb2aOTt LiNb30 s LiNbO3 LisNbO 4 -
1; 1400 2; 1600 1; 1500 2; 1750-1850 1; 1475 1; 1330 1; 1350 3; 1300-1550 6; 1475-1600 15; 1325-1700 2; 1425-1525 5; 1250-1650 2; 1500-1590 2; 1550 2; 1550-1700 2; 1550-1700
trace; 1475 pure, 1330 some, 1350 pure, 1550 pure, 1550 pure, 1550 much, 1525 some, 1500 trace, 1500 trace, 1550 trace, 1700 -
a) All compositions were sintered at 1000°C. b) The number of separate melts which were quenched at each composition. c) All glasses were colorless; amount of glass in best preparation is listed with the quenching temperature. d) For each instance the first phase to crystallize was LiNbO3.
~600 520 510 470 460 460 460 630 ~650 ~700 ~710 -
K. Nassau et aL / Ionic conductivity variation in rapitffy quenched lithium glasses
165
of necessity,involved,leastdiffusionis required when
1200 w f,. p_ 1000
/'l-iaNb04
,-LINbO 3 /LiNb308
el I--
,.I.iNbz20-t t
800
,.o..-(
/ to,,O" '
~,
6OO
-o 400
0 NblO 5
•
~
0.2
I
I 0.4
I
I
t' 0.8
0.6
J
t.0 L i 20
LI (CATION FRACTION)
Fig. 1. Phase diagram of the Nb2Os-Li20 system with circles showing the crystallization temperatures of quenched glasses. Arrows indicate compositions where glasses could
not be obtained.
.4R
t0-3
/\
/'.
T
\ \
t '5
\
10-4
b
\
i.-m
\
to O
\
IO-g
• •
10 4
0
Nbzos
\
Nb2Os-LizO b i TO0 5
I 0.2
I 0.4
I 0.6
I 0.8
1.0
Li20
ki (CATION FRACTION)
Fig. 2. Ionic conductivity variation with Li contant in quencheal glassesat 500 K for LiTaOa and the Nb2Os-Li20 system.
the composition o f the glass is the same as that of the crystallizing phase. The further away from LiNbO3, the more excess Li20 or Nb205 is present, thus implying both a longer diffusion path, and possibly a less favorable environment for diffusion. The ionic conductivity of the glasses is highest at Li = 0.513, as shown in fig. 2, and falls off rapidly in either direction. The conductivity versus 1/T curves of some of these niobate compositions, as well as of LiTaO 3 also included in fig. 2 have been published previously [2]. Quenched 6LiNbO 3 glass made from atomic weight 6 isotope of Li had the same ionic conductivity characteristics as the ordinary LiNbO3, within the experimental error. Sargeant and Roy quenched both Ta205 and Nb205 to obtain metastable high-temperature phases, but no glass was observed [8]. In general, the quenching and glass characteristics in the L i 2 0 - T a 2 0 s system of table 2 appear to be similar to those of the L i 2 0 - N b 2 0 5 system of table 1. The L i 2 0 - T a 2 0 5 phase diagram [9] shows the existence of three compounds: LiTa6016 , Li3TaO4, and LiTaO3, the latter with some solid solution similar t o that of LiNbO 3 in fig. I. Tables 1 and 2 indicate that essentially pure glass could be obtained in these systems close to the Li = 0.50 concentration. Accordingly, the rapid fall-off in both directions in fig. 2 may be partially explained by a reduction in the amount of glass present with a concentration departure from Li = 0.50 in either direction. This may not be the full explanation, since a quantity of glass giving a conductivity down 2~ orders of magnitude at Li = 0.75 would probably not be detectable by DTA. Since glass is likely to be present as a continuous phase surrounding crystallites, a large reduction in conductivity resulting from discontinuities in the electrical conductivity path would not be expected. It may be noted that the ionic conductivity of single crystal LiNbO 3 extrapolated to room temperature is ~1020 times smaller than that of quenched vitreous LiNbO 3. In the mixed LiNbO3-LiTaO 3 system glass was obtained throughout the whole system and the crystallization temperature varied monotonically with composition: from 100% LiNbO 3 to 100% LiTaO 3 there was very little observable change on the ionic conductivity. All of this indicates a smooth "isomor-
166
If. Nassau et aL / Ionic conductivity variation in rapidly quenched lithium glasses
Table 2 Quenched specimens in the Li20-Ta2Os system Li content a)
Number of quenches and temperature (°C)
Occurrence of glass b) Crystallization and temperature (*C) exotherm c) (0C)
0.33(LiTa30 a 0.50(LiTaOs) 0.60 0.80
1; 1800 7; 1685-1890 3; 1700-1900 3; 1700-1850
much, 1850 much, 1800 -
570-630 620, 650 -
a) All compositions were sintered at IO00°C. b) All glasses were colorless; amount of glass in best preparation is listed with the quenching temperature. c) For each instance the first phase to crystallize was LiTaOa. phous" substitution of Nb and Ta for each other in the vitreous phase. This is definitely not the case in the mixed LiNbO 3 : KNbO 3 : NaNbO 3 system [10] where a marked reduction of the conductivity is observed when two alkalis are present in the glass composition as shown in fig. 3. The reduction in conductivity on the mixing o f alkali ions is a phenomenon well known in the network-forming glasses such as silicates, as the "mixed alkali effect" [11]. The activation energy shows an analogous behavior [ 10]. 3.2. Variation o f ionic conductivity with lithium con ten t
In the most general form, the lithium ion conductivity o can be expressed as o = ~
i
"fui,
I0-~
Ss "q
"7
~
i
1o-4
r
'~
,
,~ ~o-~
I
)I--
I I
/
,, I
i
g
I I ~s
0'2 d , NoNbO 3
LiNbO 3
/\
% %lb. ~ w
/
0 .'6 0 .'8
KNbO 3
NoNbO s
Fig. 3. The mixed alkali effect in LiNbO s : K N b O 3 : NaNbO a q u e n c h e d glasses.
where the sum is over all types of Li ion sites i, n i is the density of ions of charge e at sites i, and #i is a mobility factor which contains the details o f the structural environment of site i. In a glass, some Li ions may be immobile with/a i = 0, while others may move freely from one site to another. Thus to maximize the ionic conductivity involves maximizing the density of mobile ions and their mobility. In general, increasing the Li ion concentration o f a glass will also change the mobility, so that the net result may not be an increase of a. In this section experiments are described which were designed to study the effects on o of varying the Li ion content in various glass systems, while attempts to vary the microstructure by chemical addition to the glass in such a way as to increase the mobility are described in the following section. The ionic conductivities at 500 K of the quenched glass systems of Li20 with Nb205, Ta205, A1203, Bi203, Ga203, WO 3 and MoO 3 are shown in fig. 4. Also included are the data from the quenched mixed sulfate glasses in the system Li2SO 4 with SrMg(SO4) 2 and La2(SO4) 3 [12,13], and from some mixed lithium glasses containing B20 3 and other ingredients as published by Smedley and AngeU [14]. The mole fractions quoted in this last work have been recalculated as cation fractions. Three characteristics can immediately be observed in fig. 4. First, there is similar behavior among equivalent valence cations which might reasonably be expected to substitute for each other, even when there is a large variation of ionic radius as that in going from the 0.53 A o f A13+ to the 1.02 A of Bi 3+. Second, there are important differences in the shapes of the curves; there are also differences in the location of
K. Nassau et aL / Ionic conductivity variaffon in rapidly quenched lithium glasses -2
.$ "7
I
I
I
/ -
I
...',.. .,/
w,i-/
/'k/
~-"
' ...~<.-"-.../
~t
Y\
Z
~--4
f
/
/",
b II -
IN "S
0
Nb,TI -7 0
I
0.2
/i
0,4 0.6 U (CATIONFRACTION)
,
0.8
1.0
Fig. 4. Ionic conductivity variation with Li content at 500 K for Li-containingoxide and sulfate glasses. glass-forming regions. Third and most surprising, the best ionic conductivities in each system are all about the same, ~10 - 3 ( ~ cm) -1 even though they occur at different Li concentrations. It is clear from these data that no simple explanation is likely to account for all the observations. The exponential increase of conductivity with lithium ion content in the case of the Li20 : A1203, Li20 : Ga203 and Li20 : Bi203 systems can be acounted for in terms of a constant ion mobility and an exponential increase of the mobile carrier concentration, even though the total lithium content only increases linearly. This explanation has been used to account for the increase of o with Na content in the SiO 2 : Na20 system [15] where the Na ions are assumed to be distributed between two types of sites: those which are associated with non-bridging oxygen ions and those which are dissociated (the weak electrolyte model). However, in the absence of any structural information in these glasses an equally plausible explanation for the exponential increase of o can be given in terms of a large number of alkali ion sites having a gaussian distribution of activation energies [4]. Within this picture it is not possible to distinguish mobile and immobile lithium ions. An approximately exponential increase of conduc-
167
tivity with lithium contant is also observed with the Li2SO 4 : La2(SO4) 3 system [12], although the exponent is only about one third of the value for the oxides just discussed. In terms of the random site mod. el this implies a narrower distribution of site energies, whereas in terms of the weak electrolyte model this would imply a smaller energy difference between associated and dissociated lithium ions. In either case this might be expected for sulfates since the negative charge on the (SO4) 2 - ion is distributed over a wider volume and would have a smaller binding energy to the lithium ions. The conductivity of the Li20 : w e 3 and Li20 : MoO 3 systems is relatively constant as the lithium ion concentration is increased. In this case, it would appear that Li ion mobili W is relatively unaffected by the composition - in agreement with studies of Li diffusion in electrochromic w e 3 films. In the Li20 : Nb205 and Li20 : Ta205 systems a pronounced maximum of the conductivity is observed for a cation fraction near 0.5. As discussed earlier, it is not likely that this is due to the formation of crystalline material within a glass since it is known that a small amount of crystallized material has little effect on the conductivity. The most reasonable explanation of the maximum is that the lithium ion mobility reaches a maximum value for 50 cation percent lithium and a further increase results in a decreased mobility. This can arise from the formation of negatively charged centers such as non-bridging oxygen ions which act as alkali ion traps in the structure. We have already pointed out tliat if the structure consists of a random array of corner linked NbO 6 octahedra, then the concentration of non-bridging oxygen ions is minimum (zero) for x = 0.5. A similar argument would predict maxima in Li20 : MgO and Li20 : T i e 2 glass systems at compositions of Li4MgO 3 and LiTiO 3 if the Mg and Ti ions are octahedrally coordinated. We have, however, been unable to prepare glasses of these compounds to verify this model. In both the Li2SO4 : SrMg(SO4) 2 glass system [13] and the mixtures reported by Smedley and AngeU [14] a decrease of the ionic conductiviW is observed with increasing Li ion concentration for high concentration values. In these cases the systems are too complex to speculate about the origin of the behavior.
168
K. Nassau et aL / Ionic conductivity variation in rapidly quenched lithium glasses
Table 3 Quenched alkali niobate and tantalate specimens with the addition of glass-formers or "vacancy" inducers Composition a)
x
LiNbOs LiNb t-xOs72.sx*xSi02 LiNbt_xOs_2.sx.xGe02 LiNbt_2xO3_sx*xB203 (1-x)LiNb0a.xLAS
0.2 0.2 0.1 0.1 0.2 0.1 0.02 0.1 0.2 0.05 0.1 0.03 0.1 0.2 0.1 0.2
(1-x)LiNbO 3 .xSLS (Lil_ 2xMgxe)x)Nb03 (Li I _ 3xAlx¢2x)Nb03
(Lil-x¢x)Nbt-xWx03 Li(Nb 1-2xWxCx) (03-2x¢2x) KNbO3 (K t - 2x Lax~2x)Nb0 3 LiTaO 3 Li(Tat_ ~xWxCx)(O3_ 2x~2x)
-
0.1 0.1 0.2
Number of quenches and temperature (°C) 15; 1325-1700 1; 1800 1; 1750 1; 1750 1; 1550 1; 1300 1; 1560 6; 1500-1800 5; 1350-1700 5; 1225-1500 1; 1550 1; 1550 4; 1475-1750 1; 1350 1; 1300 1, 1400 1; 1375 3; 1175-1310 1; 1450 7; 1685-1950 1; 1600 1; 1525
Occurrence of glass b) and exother~ (*C) pure, 460 pure, 550 c), 610 pure, 530 e), 580 much, 510 pure, 520 c) 560 pure, 620 pure, 520 c),~60L much, 460 pure, 470 some, 500 pure, 540 pure, 550 pure, 460 pure, 460 pure, 450 much, 490 some, 480 pure, 380 e), 460 pure, 620 some, 570 some, 550 some, 610
a) LAS is the lithia alhmina silicate 35 : 5 : 60; SLS is the soda lime silicate 20 : 11 : 60; ~ indicates a "vacancy". b) Amount of glass in best preparation is listed; only strong, initial crystallization exotherm is given. c) Glass transition endotherm.
3.3. The variation o f ionic conductivity with additives In the crystalline state alkali niobates and tantalates are most frequently characterized b y the corner linking o f NbO 6 octahedra in an infinite threedimensional array: for instance the perovskite (KNbO3, NaNI~O3) , tungsten bronze (Ba4Na2Nbl0030 , K6Li4Nbl0 O30), or LiNbO 3 structures are all characterized in this way. The alkali ions occupy the interstices between the NbO 6 octahedra. In the case of the perovskites there are no vacant interstitial cation sites in the structure available for the m o t i o n of alkali ions. In LLNbO 3 and the tungsten bronzes there are vacant cation sites, but the structural barriers between these sites prevent fast ion motion. The basic building blocks o f the glass are again likely to be NbO 6 octahedra, corner shared, but because of the random structure of the glass there exist paths which allow m o t i o n of alkali ions through-
out the structure which do n o t exist in the crystalline state. The question we attempted to answer was whether this ionic mobility for the glass could be increased b y introducing additional vacant cation sites b y doping with cations o f different valence. This procedure is of course well known to increase the ionic mobility in many crystalline materials. We have already shown that the addition o f MgO and A1203 to LiNbO 3 glass results in decreased conductivity, while the addition o f WO 3 produced little change o f the conductivity [3]. However, within the above model for the glass structure these additions would result in the creation o f negatively charged non-bridging oxygen ions in the glass which act as traps for the Li ions. To avoid this, KNbO 3 glass was doped with L a N b 3 0 9 , which would retain the corner linked NbO 6 framework without introduCing nonbridging oxygen ions while increasing t h e d e n s i t y o f
K. Nassau et aL / Ionic conductivity variation in rapidly quenched lithium glasses
vacant interstitial cation sites, i.e. the composition employed is most illustratively represented by
(Lil.8550.15)(Wl.85 Nb0.15)O6.85,
KO.7 Lao. 1 ~0.2 NbO 3 ,
(Lil.7~b0.15 Mg0.15)W207,
where ~ indicates a vacancy. This and other such compositions quenched are listed in table 3. The conductivity of this glass at 500 K was measured to be 2 × 10 -7 (I2 cm) -1 - about two orders of magnitude lower than previously reported for a iaure KNbO 3 glass. If the structural picture is correct, this resuli imp]i6:~'tl~at the number of vacant alkali ion sites in thepure KNbO 3 glass is already sufficient for fast ion conduction, and the reduction in conductivity is due to the reduction of the total density of alkali ions available for transport. It was shown in the preceding section how the conductivity can be an extremely non-linear function of alkali ion concentration. Of course, there may be other effects present, such as blocking of the tunnels available to K transport by the large La 3+ ions, or changes in the basic structural framework. A sirnflar set of substitutions in Li2W20 7 glass was performed since in this system the ionic conductivity is relatively insensitive to the alkali content. The glass structure is probably built upon comer linke d WO4 tetrahedra with Li ions at the interstices. We attempted to increase the density of vacant interstitial sites via the substitution of Mg2+ or A13+ for Li+, or W6+ for Nb 5+ ions. The compositions employed are shown in table 4 and are
169
(Li 1.7~0.2A10.1)W207 • Remeasurement of the Li2W207 glass indicated significant electrode polarization at 1 kHz, where previous measurements were made, and an actual conductivity at 500 K of ~-3.5 X 10- 3 ([2 cm) -1 , one order of magnitude greater than that reported from the 1 kHz measurements. This conductivity compares favorably with the best Li ion conductivities reported for glasses and also that for Li3N and Li~-alumina. All three additions to Li2W207 resulted in a decrease in the conductivity. For both the A1 and the Mg additions, a at 500 K was ~1.0 X 10 - 3 ([2 cm) -1 , and for the Nb/205 additions o at 500 K was slightly high. er, ~1.5 X 10 - 3 (~2 cm) - 1 . These decreases are small compared to that which resulted from the addition of Lal/3NbO 3 to KnbO 3. The effect on the ionic conductivity of network forming additives such as SiO2, GeO 2 and B203 to the Li20 : Nb205 and Li20 : WO 3 glass systems were studied for several compositions listed in tables 3 and 4. These additions of coruse facilitate glass formation and could have practical value if the mixed glasses had increased conductivity. In all cases, the addition of up to 20% of the glass formers produced no significant change of the ionic conductivity.
Table 4 Quenched lithium tungstate compositions with the addition of glass formers and "vacancy inducers" Composition a)
x
Number of quenches Occurrenceof glass b) and temperature CC) and exotherm (°C)
Li2W207 (1-x)Li2W2OT.xSiO 2 (1-x)Li2W2OT.xGeO2 (1-2x)Li2W2OT-2xB203 Li2W2_xO7_ 3x.xSiO2 Li2W2_xOT_ 3x.xGeO2 Li2W2_xO~_6x.xB203 (Li2_2xMgxc~x)W~07 (Li2_ 3xAlx¢2x)W207 Li2Wz_2xNb~x(OT_xCx)
0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.15 0.1 0.15
2; 925-975 1; 1350 1; 1290 1; 1300 1; 1350 1; 1250 1; 1275 1; 1210 1; 1170 1; 1180
pure, 340 c), 370 pure, 360 e), 390 pure, 420 c), 450 pure, 380 c), 410 pure, 350 e), 370 pure, 390 c), 440 pure, 370 c), 410 pure, 380 c), 400 pure, 380 e), 400 pure, 370 c), 400
a) LAS is the lithia alumina silicate 35 : 5 : 60; SLS is the soda lime silicate 20 : 11 : 60; ~ indicates a "vacancy". b) Amount of glass in best preparation is listed; only strong, initial crystallization exotherm is given. c) Glass transition endotherm.
170
IC Nassau et al. / Ionic conductivity variation in rapidly quenched lithium glasses
4. Conclusions
Acknowledgement
The most surprising result of this survey is that in each of the glass systems studied the same maximum value of the ionic conductivity, ~10 - 3 (~2 cm) -1 at 500 K was observed. However, in different glass systems the maximum occurs for different values of the lithium ion concentration indicating that no single explanation is appropriate for all glasses. Attempts to increase the ionic conductivity either by the addition of glass formers or by the addition of ions of different valence from those of the host glass were unsuccessful. These experiments indicate that the availability of vacant interstitial sites in glasses of the type studied here is not a limitation to the development of fast ion conduction. Of course, whether the atomic substitutions we have attempted resulted in vacant inter. stitial sites at all is open to question. It is dependent in part on the degree of short-range order and in part on the distance over which the concept of local charge neutrality can be expected to apply. The best way to view these glasses at present is as quenched liquids in which the liquid structure is largely preserved and, unfortunately, very little is known about the structure of these types of high melting point liquids. Concepts such as "non-bridging oxygen ions" or "vacancies" might not be appropriate in these types of solids.
We want to thank M. Grasso for assistance with many of the glass preparations and D.H. Olson for many of the conductivity measurements.
References [ 1] K. Nassau, C.A. Wang and M. Grasso, J. Am. Ceram. Soc. 62 (1979) 74, 503. [2] A.M. Glass, K. Nassau and T.J. Negran, J. AppL Phys. 49 (1978) 4808. [3] K. Nassau, M. Grasso and A.M. Glass, J. Non-Cryst. Solids 34 (1979) 425. [4] A.M. Glass and IL Nassau, J. Appl. Phys. 51 (1980) 3756. [5] IC Nassau, A.M. Glass, M. Grasso and D.H. Olson, J. Am. Ceram. Soc.127 (1980) 2743. [6] E.M. Levin and H.F. McMurdie, Phase Diagramsfor Ceramicists, 1975 Supplement (Am. Ceram. Soc., Columbus, 1975) p. 86. [7] JCPDS Powder Diffraction File, International Center for Diffraction Data, Swarthmore, PA (1977). [8] P.I. Sargeant and R. Roy. J. Am. Ceram. Soc. 50 (1967) 500. [9] F.I. Kozhaev, S.A. Shvartsman, S.I. Sadova and V.A. Sokholov, Opt. Mekh., Prom-st 45 (1978) 35. [10] A.M. Glass, K. Nassau and D.H. Olsen, in: Fast ions transport in solids, eds. P. Vashishta, J.N. Mundy and G.IC Shenoy (North-Holland, Amsterdam, 1979) p. 707. [11] J.O. Isard, J. Soc. Glass 43 (1959) 113. [12] K. Nassau and A.M. Glass, J. Non-Cryst. Solids (1981), to be published. [13] K. Nassau, A.M. Glass and M. Grasso, to be published. [14] S.I. Smedley and C.A. Angell, Solid State Commun. 27 (1978) 21. [15] D. Ravaine and J.L. Souquet, Phys. Chem. Glasses 18 (1977) 27.