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Nanoparticle Enhanced Conductivity in Organic Ionic Plastic Crystals: Space Charge versus Strain Induced Defect Mechanism Youssof Shekibi,*,† Angus Gray-Weale,| Douglas R. MacFarlane,‡ Anita J. Hill,‡,§ and Maria Forsyth*,† Department of Materials Engineering and School of Chemistry, and the ARC Centre for Electromaterials Science, Monash UniVersity, Clayton 3800, Victoria, Australia, CSIRO, PriVate bag 33, Clayton 3169, Victoria, Australia, and School of Chemistry F11, UniVersity of Sydney, New South Wales, Australia ReceiVed: February 28, 2007; In Final Form: April 27, 2007
High conductivity in solid-state electrolytes is a critical requirement for many advanced energy and other electrochemical applications. Plastic crystalline materials have shown promise in this regard, and the inclusion of nanosized inorganic particles in both amorphous and crystalline materials has indicated order of magnitude enhancements in ion transport induced by space charge or other defect enhancement. In this paper we present conductivity enhancements in the plastic crystal N,N′-ethylmethylpyrrolidinium bis(trifluoromethanesulfonyl)amide ([C2mpyr][NTf2]) induced by nanosized SiO2 particles. The addition of the nanoparticles dramatically increases plasticity and ion mobility. Positron annihilation lifetime spectroscopy (PALS) measurements indicate an increase in mean defect size and defect concentration as a result of nanoparticle inclusion. The scaling of the conductivity with size suggests that a “trivial space charge” effect is operable, although a strain induced enhancement of defects (in particular extended defects) is also likely given the observed increase in plasticity.
Introduction Plastic crystals are among the more promising solid-state materials being considered as potential electrolytes in electrochemical devices due to their high conductivity and plastic mechanical properties.1-4 The plastic phases of these materials are a type of mesophase in which the molecules are orientationally disordered, while preserving their long-range order. This orientational disorder can be explained by the potential energy barriers, which hinder the rotation of molecules, being sufficiently low to readily permit rotation and reorientation on lattice sites.3 These crystalline materials are often soft and waxy, and they are able to deform readily under relatively low forces. The use of ceramic fillers in solid electrolytes, to improve ionic conductivities in these materials, has generated interest since the first investigation of inert ceramic filler by Weston and Steel,5 and since then this avenue for enhanced ion conduction continues to be explored.6-12 Bhattacharyya and Maier13 investigated the same concept in nonaqueous liquid electrolytes and showed an enhancement in conductivity of almost an order of magnitude at the percolation threshold of the fillers. The enhancement was primarily attributed to an increased dissociation of the ionic aggregates in the surface region of the filler particles and/or the presence of an enhanced charge carrier concentration in a space charge layer surrounding the nanoparticles. The enhanced conductivity in inorganic crystal salts such as AgCl by addition of Al2O3 has been also attributed to a “space charge effect” in the interfacial region of the ceramic electrolyte near the inert nanoparticles.14 An alternative model for enhancement of ionic conductivity has also been described by Dudney15 as a model of “deformation * Corresponding authors. Fax: +61 3 99054940. E-mail: maria.forsyth@ eng.monash.edu.au (M.F.);
[email protected] (Y.S.). † Department of Materials Engineering, Monash University. ‡ School of Chemistry, Monash University. § CSIRO. | University of Sydney.
induced defects”. In this case the presence of an interface creates strain in a material which produces more defects (hence more mobility) and these defects are stabilized by the inert phase. Molecular dynamics calculations suggest that the strain induced by an interface leads to more disorder and hence can account for the increased conductivity.16 Flandin et al.17 have reported that the effect of strain on the electrical properties of composites strongly depends on the size of the filler particles and on the surface structure. The enhancement of conductivity in yttriastabilized zirconia (YSZ) with the increase of plastic strain has also been reported by Otsuka et al.18 Thus it appears that there are two possible explanations to account for the enhanced conductivity in nanocomposite electrolyte materials.16,19-22 In the present work we have extended the space charge and strain induced defect concepts to other soft matter electrolytes, specifically organic ionic plastic crystalline materials. These materials offer a promising alternative to polymer electrolytes as soft matter solid-state materials for devices including solar cells, batteries, and fuel cells.23-25 This paper reports a 2 order of magnitude increase in conductivity in [C2mpyr][NTf2] as a result of incorporation of nanoparticulate SiO2 and the dependence of the effect on size of the nanoparticle as probed by scanning electron microscopy (SEM), proton nuclear magnetic resonance (1H NMR) spectroscopy, and positron annihilation lifetime spectroscopy (PALS). Experimental Methods Sample Preparation. The synthesis of N-methyl-N-ethylpyrrolidinium bis(trifluoromethanesulfonyl)amide ([C2mpyr][NTf2]) has been described elsewhere.26 Samples were prepared by mixing the [C2mpyr][NTf2] and SiO2 nanoparticles above the
10.1021/jp071631j CCC: $37.00 © 2007 American Chemical Society Published on Web 07/07/2007
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TABLE 1: Entropy for the Different Phase Transitions as a Function of Filler Size as Calculated from the Peak Area of the DSC Traces and the Transition Temperature wt % SiO2
IV f III ∆S (J K-1 mol-1), (10%
III f II ∆S (J K-1 mol-1), (10%
II f I ∆S (J K-1 mol-1), (10%
I f melt ∆S (J K-1 mol-1), (10%
0 10 (7 nm) 10 (12 nm) 10 (16 nm)
4.7 1 4 1.3
3.6 2.4 3.2 2.8
2.2 1.8 2 1.8
26 24 26 25
melting point of the plastic crystal for at least 12 h using a shearing mixer technique (SiO2 was used as received from Degussa, types aerosil 300, 200 and 130, having primary particle sizes of 7, 12, and 16 nm, respectively). DSC Measurements. A TA Instruments Model Q100 differential scanning calorimeter (DSC) was used to measure the thermal properties of the samples, over a temperature range of -150 to 140 °C with a scanning rate of 10 °C min-1. Three consecutive runs were done for each sample, and the data reported in Table 1 are the average values of these three runs. The DSC traces were analyzed with the TA Instruments universal analysis 2000 program. The entropies presented are based on the amount of plastic crystal in each sample, i.e., subtracting the amount of filler in the calculations. Conductivity Measurements. The ionic conductivity of all samples was evaluated using ac impedance spectroscopy. The measurements were performed with a frequency response analyzer (FRA, Solartron 1296), driven by Solartron impedance measurements software version 3.2.0. The samples were pressed into pellets approximately 2-3 mm thick, which were then sandwiched between two steel electrodes. Samples were handled, and the cells were sealed, within a nitrogen-filled drybox. The data were collected over a frequency range of 1 MHz to 0.01 Hz and a signal voltage of 0.1 V in a temperature range from -50 to 80 °C at 5 and 10 deg intervals. The temperature was controlled using a Eurotherm Model 2204 temperature controller. The PID parameters set on the controller led to 15 min intervals between measurement temperatures (i.e., for a 5 or 10 °C ramp, the ramp rate was approximately 0.4 or 0.8 °C/ min, respectively). A short equilibration time of up to 2 min stabilized the temperature before the measurement began. (The measurement itself takes approximately 12 min for a frequency range of 106-0.01 Hz.) Multiple conductivity measurements were made at some temperatures to check for stabilization and reproducibility in these samples, and it was found that the conductivity was relatively stable during the time frame of the measurements (a slight increase of 10% by the third run). The data typically presented a single semicircle from which the conductance of the samples was determined using the real axis intercept in the Nyquist plot of the impedance data. NMR Studies. The 1H and 19F nuclear magnetic resonance spectral information and line width measurements, at Larmor frequencies of 300 and 283.2 MHz, respectively, were obtained on pure [C2mpyr][NTf2], as well as on the nanocomposite samples, as a function of temperature from room temperature up to ∼80 °C, which is just below the melting point of the plastic crystal. Results for higher temperatures have been reported elsewhere for pure [C2mpyr][NTf2].4 The typical pulse length was 13 µs with 16 scans for each spectrum. PALS Experiment. Positron annihilation lifetime spectroscopy measurements as a function of temperature were performed using a thermally stabilized automated EG&G Ortec fast-fast coincidence system under dry nitrogen atmosphere. A Mylar NaCl spot source was sandwiched between two identical pellet samples, for size and weight, with a thickness of 1.1 mm. The measurements were performed on pure [C2mpyr][NTf2] and on
samples containing 10 wt % SiO2. The lifetime spectra of the plastic crystals were resolved into two components (the shortest, in the picosecond region, is attributed to free and trapped positrons, and the longer one, τops, is attributed to orthopositronium pick-off annihilation) using the program PFPOSFIT. Contributions from the annihilations in the filler are included in the shorter values. Only the orthopositronium annihilation parameters (τops, Iops) are reported here as they are assumed to relate to the defect volume of the plastic crystal.27 SEM Imaging. Ambient temperature scanning electron microscopy was performed with a Philips XL30 field emission gun (FEG) SEM. Samples were transferred into the microscope observation chamber through a Polaron LT7400 cryoprep stage and placed under a vacuum of 3.7 × 10-6 mbar. Images were acquired with a secondary detector to reveal detailed surface topography. Low accelerating voltages of 2 kV for pure sample and up to 10 kV for samples with silica were used to minimize sample charging and damage. Each sample was prepared as a pellet and was fixed in a sample holder and then fractured at room temperature (below the II-I transition). Results and Discussion Thermal and Conducting Properties. Figure 1 shows the thermal traces of the samples of pure [C2mpyr][NTf2] and composite samples with 10 wt % SiO2 nanoparticles. The DSC traces are normalized with respect to the amount of plastic crystal in each sample. The pure material goes through three solid-solid phase transitions before melting at ∼90 °C, as reported.28 It can also be seen that the melt transition is skewed for the composite samples and the onset temperature is shifted toward lower temperatures, while the entropy change seems to be slightly affected by the filler content.
Figure 1. DSC traces of pure [C2mpyr][NTf2] and nanocomposite samples with 10 wt % SiO2 nanoparticles of 7, 12, and 16 nm.
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Figure 2. Conductivity as a function of temperature of pure [C2mpyr][NTf2] and [C2mpyr][NTf2] combined with 10 wt % silica nanoparticles of 7, 12, and 16 nm.
Figure 4. (a) 19F NMR line width measurement on pure [C2mpyr][NTf2] and [C2mpyr][NTf2] with added 10 wt % SiO2. (b) 19F MAS spectrum at 298 K spinning at 8 kHz.
Figure 3. 1H NMR line width spectra for [C2mpyr][NTf2] and [C2mpyr][NTf2]/10 wt % SiO2 nanocomposite at temperatures before and after the solid-solid phase transition from phase II to phase I.
We present the conductivity of pure [C2mpyr][NTf2] and [C2mpyr][NTf2] combined with 10 wt % silica nanoparticles in Figure 2. The conductivity measurements for all three composite samples show Arrhenius behavior across the entire temperature range encompassing phases II and I with no apparent steps in conductivity at the phase II f phase I solid-solid phase transition (see DSC traces). There is a notable 2 order of magnitude increase in conductivity with the addition of nanoparticles to this plastic crystalline material. For example, the conductivity at 60 °C (which is in the plastic phase, phase I) is 10-7 S cm-1 for the pure material and 10-5 S cm-1 for the nanocomposite sample. The gradual change of ionic conductivity in phase II also suggests that some degree of disorder persists in this lower temperature solid phase, promoting the translational motion of some fraction of ions. In several papers, as in this work, it has been observed that the particle size has an important influence on the enhancement of ionic conductivity, where the smaller size offers the larger enhancement.11,29 NMR Analysis. Figure 3 displays 1H NMR spectra at selected temperatures for the pure [C2mpyr][NTf2] and for the samples with different particle size SiO2. The spectra of the composite systems exhibit a narrower full width at half-maximum (fwhm) line width at all temperatures compared with the pure material; for example, the fwhm decreases from 29 to 17 kHz at room temperature. A further reduction in the line widths of all three composite samples was observed as the temperature increased, indicating an increase in cation rotational mobility which averages the dipole-dipole interactions that lead to broad NMR lines. The appearance of a narrow resonance superimposed on the broad peak is also evident in the nanocomposite systems.
Figure 5. Orthopositronium lifetime, τops, and orthopositronium intensity, Iops,of pure [C2mpyr][NTf2] and [C2mpyr][NTf2] combined with 10 wt % SiO2 as a function of temperature. The error in all the data points is calculated from standard deviation of the PALS fits and is individual for the pure sample.
This narrow peak has an fwhm less than 4 kHz and is likely to be associated with those ions undergoing translational motion.30 Its intensity grows with higher temperatures, indicating an increase in the number of diffusing cations, which would account for an increased conductivity as observed in Figure 2. A reduction in the 19F line width of samples with SiO2 was also observed as the temperature increased, indicating an increase in anion mobility. The 19F spectra of the [C2mpyr][NTf2] with 10 wt % SiO2 (Figure 4) exhibit smaller line widths at all temperatures compared with that of the pure [C2mpyr][NTf2] and show the appearance of a narrower peak superimposed on the broad peak. This indicates that both anion mobility and cation mobility are likely to increase in the nanocomposite as compared to the pure material. Defects and Plasticity. Figure 5 shows the orthopositronium lifetime, τops, and related intensity, Iops, of pure [C2mpyr][NTf2] and the filled samples with 10 wt % silica. The result shows an increase in τops, which is related to defect size27 compared with the pure plastic crystal over the whole temperature range studied.
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Figure 6. SEM fractured images (a) of pure [C2mpyr][NTf2] and (b) of composite sample with 10 wt % SiO2 at room temperature.
A positive deviation Iops, related to vacancy concentration, compared to the pure plastic crystal is also observed. It is interesting that this increase in both PALS lifetime and intensity (i.e., defect size and number) is equivalent to the defect expansion due to a 50 °C increase in temperature in the pure [C2mpyr][NTf2] crystal. Thus, the incorporation of nanosized SiO2 has the same effect on defects (and thereby ion mobility) as increasing temperature. Figure 6a shows the grain boundaries existing in bulk [C2mpyr][NTf2], as well as slip planes, which are dependent on the inter- and intramolecular bonding interactions and contribute plasticity. These latter features occur along a single slip direction and terminate at grain boundaries while retaining their coherency. The grain boundaries for samples with added SiO2 are very difficult to detect; however, the appearance of dramatic
plastic flow is observed in this case (Figure 6b). Given that the fracture occurred at room temperature in phase II (not the highest temperature phase), this suggests a higher level of ionic/ molecular mobility in this phase. This is consistent with an increase in defect volume (size and number) as well as the higher conductivities seen in Figure 2 and the slight increase in the narrow component observed in the 1H NMR measurements. Conduction Mechanism. In order to investigate whether the enhanced conductivity could result from a space charge region as described by Maier,14 some length scales and material properties including the interparticle distances l, estimated layer thickness around each particle 1/a, surface area density a, and number density of particles ν were calculated from the different
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TABLE 2: Length Scales in Nanoparticle/Plastic Crystal Electrolytes with Three Different SiO2 Particle Sizes r (nm)
a (nm-1)
7 12 16
2.6 × 1.5 × 10-2 1.1 × 10-2 10-2
1/a (nm)
ν (nm-3)
38.9 66.7 88.9
4.2 × 10 8.3 × 10-6 3.5 × 10-6 -5
SiO2 particle sizes (Table 2). The apparent Debye length was also calculated from eq 1:14
λ)
x
τ0kBT 2e2c∞
(1)
l (nm)
F (nm-3)
F-1/3 (nm)
15 25 34
6 × 10 6 × 10-2 6 × 10-2
2.5 2.5 2.5
-2
It is not possible to deduce the nature of the activated process from the activation energies alone or from comparison with much simpler materials, but at least there is no evidence of a change in the mechanism of charge transport with the addition of nanoparticles. The similar activation energies in pure and
In this way we can compare the two proposed space charge effects discussed by Maier et al., i.e., “trivial” effects and the “true” space charge region. A trivial nanoscale effect is one that scales with the surface area of the interfaces present. The scaling is trivial, not the effect, and “trivial space charge” effects can be very large. The “true space charge” appears, for example, when the space charge layers of neighboring interfaces overlap, so that the concentration of charge carriers becomes greater and nonlinear effects become important. This is a “true” nanoscale effect in that it is not merely the consequence of scaling a macroscopic phenomenon down to nanometer sizes; rather, it has no macroscopic analogue. If the conductivity enhancement in [C2mpyr][NTf2] is a simple effect as in layered nanostructures, we might have (see ref 13 for a similar treatment)
σ ) βσ0(1 - φ) + K∆σλa
(2)
where σ0 is the conductivity of the pure material; β and K are dimensionless factors accounting for percolation, close to unity if most pathways percolate;13 φ is the volume fraction of nanoparticles, about 6%; ∆σ is the increase in conductivity in a layer of thickness λ at the silica surface; a is the silica surface area per unit volume. This equation says that the pure material and the layers of enhanced conductivity are essentially connected in parallel. The interparticle distances are well within the range suggested previously in a CaF2/BaF2 system, so we could argue that a “true” space charge region exists, thus accounting for the increased conductivity.19 However, when the Debye length is calculated, assuming the number of charge carriers obtained from NMR measurements, it is found to be in the range of several angstroms, which is at least an order of magnitude smaller than the interparticle distances. Hence, we would argue that this interface is unlikely to be in a “true” space charge regime. In this scenario the enhanced ionic conductivity could only be attributed to a “trivial” space charge effect in the interfacial region of the electrolyte near the nanoparticle. Figure 7 shows that the conductivity is roughly proportional to the silica surface area density at each temperature. The 16 nm data (a ) 0.011 nm-1) seem systematically a little low, the 12 nm data (a ) 0.015 nm-1) are perhaps a little high, but the conductivity data are consistent with this trivial scaling. Without clear evidence for a nontrivial (“true nanoscale”) effect, as in ref 31, it is reasonable to assume that eq 2 holds. Furthermore, Figure 8 shows an Arrhenius plot of the four conductivities, along with K∆σλ, the coefficient of a in fits of eq 2 to the data in Figure 7. It allows the extraction of the activation energies which are close to 90 kJ mol-1 (see Table 3). This value is about 1 eV, typical of the activation energies of point defect mediated conduction in ordinary superionic crystals such as lead(II) fluoride and calcium fluoride.
Figure 7. Electrical conductivities for pure [C2mpyr][NTf2] and [C2mpyr][NTf2] with added silica nanoparticles against silica surface area density. The lines are straight-line fits at each temperature. The linear dependence suggests that most available paths percolate. The two plots differ only in vertical scale.
Figure 8. Arrhenius plot of four conductivities for the four samples, along with K∆σλ, the coefficient of a in fits of eq 2 to the data in this figure.
11468 J. Phys. Chem. C, Vol. 111, No. 30, 2007 TABLE 3: Activation Energies for Conductivities and for the Area Density Coefficient of Conductivity Increase (See Eq 2) -1
quantity
EA (kJ mol )
EA (hartree)
pure σkT 7 nm σkT 12 nm σkT 16 nm σkT K∆σkT
95 86 92 91 92
0.036 0.033 0.035 0.035 0.035
nanostructured material argue that the plastic crystal has the same basic structure in the two materials, though more disordered in the nanostructured case. Combined with the apparent trivial scaling, this result implies that some straightforward interfacial layer effect is responsible for the conductivity increase. The best information about the thickness of the layer comes from the fact that the data are consistent with eq 2. This equation assumes that the layers around the nanoparticles are connected in parallel with the remaining volume of unaffected plastic crystal. This is a reasonable approximation, and we can assume that the factors β e 1 and K e1 account for the errors so introduced.13 If the layer of enhanced mobility around each nanoparticle were much thinner than the interparticle spacing, then the conductivity would not be proportional to the silica surface area density because it would be dominated by the regions of low, pure-crystal conductivity. On the other hand, if the thickness were much greater than the interparticle spacing, then the enhancement would saturate, as more surface area would not increase the volume of affected crystal. It seems fair to assume that the layer thickness is λ ∼ 10 nm, provided we are careful that this is an order of magnitude estimate only; λ ) 1 nm and λ ) 100 nm are very unlikely (see Table 2 for interparticle distances). On the other hand, the SEM images show dramatically higher plasticity with ductile fracture surfaces obtained in the nanocomposite materials, even when the fracture occurs in phase II. Comparing with the smoother appearance of the fracture surface obtained in the case of the pure [C2mpyr][NTf2], this would suggest that the presence of so much surface area (SiO2) provides sufficient strain in these soft materials to increase the number of defects, thus leading to plastic flow and to ductile failure. Additionally, the 1H NMR data indicated a higher level of cation mobility with the addition of these particles in the entire range of temperature and the PALS data provided evidence of an increase in both defect size and number on addition of smaller nanosize SiO2. These observations are consistent with an increased conductivity created via “strain induced defects” as proposed by several researchers.15,16,21,22 The soft nature of the materials in the series of organic ionic plastic crystals should increase this possibility. Onuki has described and studied a time-dependent Ginzburg-Landau theory of plastic crystals, neglecting any electrostatics, and reports the response of vacancy concentrations to plastic strain.32 This effect will be further investigated in future work.
Shekibi et al. Conclusion A significant increase in conductivity induced by nanoparticle SiO2 in these soft electrolyte materials appears to be associated with the increased number of mobile defects as evidenced by PALS and solid-state NMR coupled with increased plastic flow observed in the SEM images. The analysis presented here points to a strain induced defect mechanism rather than a true space charge mechanism. If this nanoparticle effect is also observed in doped plastic crystal systems, then it may be possible to produce target ion diffusivities (e.g., Li+, H+) which will lead to higher performance in all solid-state lithium batteries and fuel cell devices. This is currently under investigation in this and related plastic crystal systems. References and Notes (1) Alarco, P. J.; Abu-Lebdeh, Y.; Ravet, N.; Armand, M. Solid State Ionics 2004, 172, 53. (2) Wang, P.; Dai, Q.; Zakeeruddin, S. M.; Forsyth, M.; MacFarlane, D. R.; Gra¨tzel, M. J. Am. Chem. Soc. 2004, 126, 13590-13591. (3) MacFarlane, D. R.; Forsyth, M. AdV. Mater. 2001, 13, 957. (4) MacFarlane, D. R.; Huang, J. H.; Forsyth, M. Nature 1999, 402, 792. (5) Weston, J. E.; Steel, B. C. H. Solid State Ionics 1982, 7, 75. (6) Chung, S. H.; Wang, Y.; Persi, L.; Croce, F.; Greenbaum, S. G.; Scrosati, B.; Plichta, E. J. Power Sources 2001, 644, 97-98. (7) Capuano, F.; Croce, F.; Scrosati, B. J. Electrochem. Soc. 1991, 138, 1918. (8) Borghini, M. C.; Maslmgostino, M.; Passerini, S.; Scrosati, B. J. Electrochem. Soc. 1995, 142, 2118. (9) Croce, F.; Appetecchi, G. B.; Scrosati, L. P. B. Nature 1998, 394, 456. (10) MacCallum, J. R.; Seth, S. Eur. Polym. J. 2000, 36, 2337. (11) Adebahr, J.; Best, A. S.; Byrne, N.; Jacobsson, P.; MacFarlane, D. R.; Forsyth, M. Electrochim. Acta 2003, 48, 2099. (12) Adebahr, J.; Best, A. S.; Byrne, N.; Jacobsson, P.; MacFarlane, D. R.; Forsyth, M. Phys. Chem. Chem. Phys. 2003, 5, 720. (13) Bhattacharyya, A. J.; Maier, J. AdV. Mater. 2004, 16, 811. (14) Maier, J. Prog. Solid State Chem. 1995, 23, 171. (15) Dudney, N. Solid State Ionics 1988, 28-30, 1065. (16) Gray-Weale, A.; Madden, P. A. J. Phys. Chem. B 2004, 108, 6634. (17) Flandin, L.; Chang, A.; Nazarenko, S.; Hiltner, A.; Baer, E. J. Appl. Polym. Sci. 2000, 76, 894. (18) Otsuka, K.; Kuwabara, A.; Nakamura, A.; Yamamoto, T.; Matsunaga, K.; Ikuhara, Y. Appl. Phys. Lett. 2003, 82, 877. (19) Maier, J. Solid State Ionics 2003, 157, 327. (20) Lupascu, D. C. Phys. ReV. B 2004, 70, 184124. (21) Gray-Weale, A. R. Soc. Chem., Faraday Discuss. 2006, 134, 17. (22) Su, J.; Zhang, Q. M.; Ting, R. Y. Appl. Phys. Lett. 1997, 71 (3), 386. (23) Dai, Q.; MacFarlane, D. R.; Forsyth, M. Solid State Ionics 2006, 177, 395. (24) Long, S.; MacFarlane, D. R.; Forsyth, M. Solid State Ionics 2004, 175, 733-738. (25) Abu-Lebdeh, Y.; Abouimrane, A.; Alarco, P.-J.; Dividson, I.; Armand, M. J. Power Sources 2006, 159, 891-893. (26) MacFarlane, D. R.; Meakin, P.; Sun, J.; Amini, N.; Forsyth, M. J. Phys. Chem. B 1999, 103, 4164. (27) Pas, S. J.; Huang, J. H.; Forsyth, M.; MacFarlane, D. R.; Hill, A. J. J. Chem. Phys. 2005, 122, 064704. (28) Adebahr, J.; Ciccosillo, N.; Shekibi, Y.; MacFarlane, D. R.; Hill, A. J.; Forsyth, M. Solid State Ionics 2006, 177, 827-831. (29) Krawiec, W., Jr.; Scanlon, L. G.; Feller, J. P.; Vaia, R. A.; Vasudevan, S.; Giannelis, E. P. J. Power Sources 1995, 54, 310. (30) Adebahr, J.; Seeber, A. J.; MacFarlane, D. R.; Forsyth, M. J. Appl. Phys. 2005, 97, 093904.1-093904.5. (31) Sata, N.; Eberman, K.; Eberl, K.; Maier, J. Nature 2000, 408, 946. (32) Onuki, A. J. Phys.: Condens. Matter. 2003, 15, S891.
