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Li2S–Li2O–P2S5 solid electrolyte for all-solid-state lithium batteries
Solid State Ionics 214 (2012) 25–30
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Solid State Ionics journal homepage: www.elsevier.com/locate/ssi
Li2S–Li2O–P2S5 solid electrolyte for all-solid-state lithium batteries James E. Trevey ⁎, Jeremy R. Gilsdorf, Sean W. Miller, Se-Hee Lee Department of Mechanical Engineering, University of Colorado-Boulder, Boulder, CO 80309-0427, USA
a r t i c l e
i n f o
Article history: Received 2 November 2011 Received in revised form 17 February 2012 Accepted 20 February 2012 Available online 28 March 2012 Keywords: Solid-state Lithium Battery Sulfur Lithium cobalt oxide
a b s t r a c t Various compositions of Li2S–Li2O–P2S5 solid electrolytes were prepared by planetary ball milling and systematically investigated. Structural analysis by X-ray diffraction (XRD) and Raman spectroscopy showed phase changes in the electrolyte corresponding to changes in conductivity. All-solid-state LiCoO2/Li cells were tested by constant–current–constant–voltage (CCCV) charge–discharge cycling at a current density of 50 μA cm− 2 between 2.5 and 4.3 V (vs. Li/Li+). In spite of reduced conductivity from introducing oxide species into a sulfide based electrolyte, increased cycle stability is observed for increased concentrations of Li2O in the electrolyte. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Advancements in energy storage technology for applications ranging from portable electronics to solar energy are continually sought by modern science. Lithium-ion batteries have always attracted considerable attention as rechargeable power sources due to high achievable energy and power densities, as well as good shelf life and no memory effect [1–7]. However, typical lithium-ion batteries containing liquidelectrolyte maintain numerous safety issues regarding flammability and toxicity [8–14]. These safety issues not only require considerable “dead weight” in regards to the devices and insulation required for controlling and containing the hazardous liquid which reduces energy density, but restrict scale up for potential applications regarding PHEV and EV due to the higher risk associated with increased size of liquidelectrolyte batteries [1,6,15–20]. Therefore improving safety of rechargeable lithium-ion batteries has become of great interest to the research community with respect to overall safety and performance [10,21]. This drives research for improvement of all-solid-state batteries which are inherently safer, as they do not contain any liquids, exhibit leakage, or have potential to undergo thermal runaway [5,8,12,14,18,20,22]. Though currently, all-solid-state batteries cannot match the performance of liquid-based batteries due to low achievable ionic conductivity and interfacial instability of solid electrolytes with electrode materials [4,10,13,23,24]. A large variety of solid electrolytes have been identified for use in all-solid-state batteries with ranging properties and target characteristics. The sulfide based electrolyte Li2S–P2S5 has been of continued ⁎ Corresponding author at: HRL Laboratories LLC, Malibu, CA 90265-4797, USA. Tel.: +1 310 317 5137; fax: +1 310 317 5840. E-mail address: [email protected] (J.E. Trevey). 0167-2738/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2012.02.034
research for over two decades for its high ionic conductivity, stability in contact with lithium metal, and simplicity in design [3,11,15,17,24–27]. Numerous additives to the electrolyte such as GeS2, GeSe2, P2O5, P2S3, SiS2 and more have been introduced to the electrolyte to aid in interfacial stabilization of solid electrolytes with positive electrode materials such as LiCoO2 [2,6,23,28]. However, research has shown limited success in stabilizing the capacity fade of solid batteries with sulfide based electrolytes and oxide based positive electrode materials. The difference in potential between these species causes the development of highly resistive reaction products within batteries that results in rapid capacity fade [5,10–12,14,20–22]. Opposed to melt and quench techniques that have been employed in the past; the use of ball milling as a means of producing solid electrolytes offers numerous advantages [11,25]. From the standpoint of manufacturability, ball milling is lower cost and less time consuming, while from a materials science aspect, ball milling is capable of producing ultra-fine powders sufficient for high conductivity of solid electrolyte and good interfacial contact between electrolyte and electrode materials [14,20]. Ball milling embraces a complex mixture of fracturing, grinding, high-speed plastic deformation, cold welding, thermal shock, and intimate mixing [26–30]. The mechanical energy created by the milling can induce chemical reactions of the reactants at room temperature [31], introduce a large number of crystal defects, and increase the specific surface area of compounds, thought to be one of the main reasons for increased reactivity of milled materials [32,33]. The present paper reports on the ternary solid electrolyte system Li2S–Li2O–P2S5 generated by planetary ball milling regarding characterization and application to all-solid-state secondary batteries. In an effort to reduce the potential between the sulfide based electrolyte and oxide based cathode, LiCoO2, an oxygen containing species, Li2O, was introduced to the electrolyte. In concept, use of the ternary
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electrolyte having sulfide and oxide lithium-containing species will result in reduced side reactions and improved cycling stability of batteries. In order to get a global view of the system, three sets of data with varying compositions were manufactured, characterized, and tested. With the typical range of x = 70–80 in the binary xLi2S-(100x)P2S5 system for highly conducting and stable electrolyte, we used a similar compositional scheme in which the molar amount of P2S5 was fixed, and the lithium containing compound ratios were varied. Our results detail an investigation of the ternary electrolyte in the ranges of xLi2S-(80-x)Li2O–20P2S5 for x = 0–80 (−20P2S5), xLi2S-(75-x)Li2O–25P2S5 for x = 0–75 (−25P2S5), and xLi2S-(70-x) Li2O–30P2S5 for x = 0–70 (−30P2S5). 2. Experimenatal Li2S–Li2O–P2S5 solid electrolytes (SE) were prepared by planetary ball milling. Reagent Grade Powders of Li2S (Alderich, 99.999%), P2S5 (Alderich 99%) and Li2O (Alfa-Aesar 99.9%) were used as starting materials. Appropriate concentrations of materials were combined into a stainless steel agate vial. SE was prepared using 2 large stainless steel balls (10 mm diameter) and 6 small stainless steel balls (6 mm diameter) for grinding. High energy ball milling (Across International PQ-N4) took place for 20 continuous hours in a sealed Ar container. Composite electrodes were prepared by mixing LiCoO2 powder (Sigma Alderich), SE Li2S–Li2O–P2S5 and acetylene black (Alfa-Aesar, 50% compressed) at a weight ratio of 20:30:3 respectively. Bilayer electrolyte pellets are formed by hand pressing 100 mg of Li2S–Li2O– P2S5 on top of a 100 mg hand pressed layer of 77.5Li2S-22.5P2S5 (mol%) SE. A 10 mg layer of the composite material is then carefully spread on the top of the Li2S–Li2O–P2S5 electrolyte layer and the cell pelletized by cold pressing (5 metric tons) for 5 min. Li foil (AlfaAesar, 0.75 mm thick) is then attached to the 77.5Li2-22.5P2S5 (mol%) SE face that was previously shown to be stable against Li metal [34], at 2 metric tons. All pressing and testing operations are carried out in a polyarytherketone (PEEK) mold (ϕ = 1.3 cm) with Ti metal rods as current collectors for both working and counter electrodes. All processes were carried out in an Ar-filled glove box. Galvanostatic charge–discharge cycling took place at first cycle cut off voltages of 4.3 and 2.5 V at a current of 50 μA cm − 2 at room temperature using an Arbin BT2000. SE samples were characterized by XRD measurement with Cu-Kα radiation. Samples were sealed in an airtight aluminum container with Beryillium windows and mounted on the X-ray diffractometer (PANalytical, PW3830). Raman spectroscopic measurments were taken through sealed glass containers (Jasco NRS-3100). Ionic conductivities were measured by AC impedance spectroscopy (Solarton 1280 C) of which weighed materials are cold pressed to both sides
of the SE pellet to serve as electrodes. The impedance of selected cells was measured from 20 MHz to 100 mHz at room temperature and the conductivity was determined using complex impedance analysis. Schematic diagrams for the Li/SE/LiCoO2 cells and AC impedance cells can be seen Fig. 1(a) and (b), respectively. 3. Results Although sulfide based solid electrolyte (thio-LISICONs) are preferable over oxide based solid electrolytes (LISICONs) for their high polarizability and large ionic radius that allow for higher conductivity, sulfide based electrolytes are not completely stable against oxide based cathode materials such as LiCoO2 due to the large potential difference. While the introduction of an oxide-based species, Li2O, into the sulfide based solid electrolyte is expected to cause a loss of conductivity, it is proposed here that incorporation of Li2O into the binary Li-P-S electrolyte will mitigate the potential difference between the oxide-based cathode and sulfide based electrolyte, thereby reducing any side reactions that may result in increased internal resistance within the cell. Fig. 2 shows impedance data for two cells made with different solid electrolytes, one rich in Li2O and one poor in Li2O. The data depict the impedance measurements immediately after cell fabrication, after first charge, and after first discharge. As expected, higher overall impedance is observed for the sample with more Li2O. However, the normalized difference from the chargeinterfacial impedance and the discharge-interfacial impedance of the two samples shows a 6% reduction of interfacial impedance for the Li2O rich cell. Increased stability of the electrode/electrolyte contact is projected to outweigh the conductivity losses of the solid electrolyte. Fig. 3 shows a comparison of the conductivities of the three ball milled electrolyte systems. A generally decreasing conductivity can be observed for decreasing x, or decreasing Li2S content and increasing Li2O content, as expected. Furthermore, with an increased molecular composition of P2S5, a similar trend of decreasing conductivity can be observed, for instance, with the maximum conductivity of the –20P2S5, –25P2S5, and the –30P2S5 systems being 3.0 × 10− 4 S cm- 1, 1.6 × 10− 4 S cm- 1, and 5.4 × 10− 5 S cm − 1 respectively, resulting from smaller concentrations of Li mobile ions in the electrolyte [35,36]. This trend agrees with our previous work regarding single step ball milling (SSBM), but with lower overall conductivity [36]. Machida et al. also tested the Li2S–Li2O–P2S5 system and found a high conductivity of 2.7 × 10− 4 S cm− 1 for a composition of 67.5Li2S–7.5Li2O–25P2S5, of which composition we calculate a conductivity of 1.1 × 10− 4 S cm− 1 for our electrolyte [37]. The difference in conductivity values is relatively small and likely due to minute differences in milling conditions. While the decreasing conductivity of the –20P2S5 system with decreasing x remains relatively consistent, the –25P2S5 system shows a region
Fig. 1. Schematic diagram of Ti test die for a) the all-solid-state battery and b) measuring conductivity of solid-electrolyte.
J.E. Trevey et al. / Solid State Ionics 214 (2012) 25–30
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are used for standardization of the data. Throughout all 3 groups of electrolyte, highly crystalline peaks for Li2S at 26.8° are observed at high values of x, while peaks at 34° for Li2O are observed for lower values of x [38,39]. A broad amorphous peak is observed for all 3 sample groups from 25° to 35°, with the –20P2S5 samples showing a second broad peak at 17°–18° that briefly appears for high values of x in the – 25P2S5 patterns as well. While planetary ball milling typically results in amorphous electrolyte, our electrolyte exhibits some crystalline behavior as a result of the increased temperature of the ball milling vials during the 20 consecutive hours of rotation, similar to the SSBM formation of highly crystalline and highly conducting materials by ball milling in a high temperature environment [34].
Fig. 2. Impedance for (a) x = 50 and (b) x = 75 of the xLi2S-(80-x)Li2O–20P2S5 electrolyte system.
of markedly lower conductivity between x = 35–50, and the –30P2S5 system shows a rapid decrease in conductivity as x approaches zero with the exception of one point at x = 20. For both the –20P2S5 and the –25P2S5 systems there is approximately 1 order of magnitude difference in conductivity between x = maximum, and x = 0 while the –30P2S5 system shows a decrease in conductivity of 6 orders of magnitude. X-ray diffraction patterns for the electrolytes can be observed in Fig. 4 with significant differences between the three systems. Shown by group according to the molar amount of P2S5, a distinct decrease in relative intensity is observed for increasing P2S5 content. Small peaks at 25° and 38° are consistent with records for Be and
Fig. 3. Conductivity map of the –20P2S5, –25P2S5, and –30P2S5 solid electrolyte systems.
Fig. 4. X-ray diffraction comparison of the –20P2S5, –25P2S5, and –30P2S5 solid electrolyte systems for increasing value of x by increments of 5.
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Interestingly, comparison between the XRD patterns and the conductivity map from Fig. 3 exhibits significant correlation between data sets. In regards to the –25P2S5 system, a substantial intensity and crystallographic difference for the region x = 35–50 as was observed by a marked decrease in conductivity for the same range of x in Fig. 3. In this data range, the development of a broad peak around 16° which has been previously attributed to a thio-LISICON III analog [40] is observed, as well as a highly crystalline peak corresponding to the presence of Li2O. Based on the results of XRD, the drop in conductivity of the –25P2S5 system is a result of increased Li2O presence in the electrolyte. Similarly for the –30P2S5 system, at the composition x = 20, a sharp increase of conductivity was observed in Fig. 3, which correlates to the disappearance of the Li2O peak in the corresponding XRD pattern. The Raman spectra of Li3PS4 and Li4P2S6 crystals have been previously displayed by Machida et al. [41]. The Li3PS4 spectrum exhibits a strong peak at 424 cm − 1 and two weak peaks at 550 and 567 cm − 1 of which Tachez et al. [42] ascribed to tetrahedral PS43- unit existing in the Li3PS4 crystal. The Li4P2S6 spectrum showed a strong peak at 387 cm − 1 and weak peaks at 550, 570, and 257 cm− 1 of which Mercier et al. [43,44] reported that such peaks can be attributed to PS3 terminal modes of P2S64 − unit; i.e., the P–P bonded hexathiohypophosphate group. In Fig. 5, Raman spectroscopic measurement of all three systems of materials reveals a number of different relationships between material compositions, X-ray diffraction patterns, and conductivity measurements. The –20P2S5 system predominantly showed evidence of the Li3PS4crystal, the intensity of which increases with increasing composition of Li2S, indicating that increased Li2S content leads to the formation of the PS43- unit and higher conductivity. The –25P2S5 system exhibited characteristics of the Li3PS4 crystal, as well an unknown peak near 406 cm - 1 at the midpoint between Li3PS4and Li4P2S6 crystals. A sharp shift from of the strong peak is observed near x = 35, similarly to X-ray diffraction in Fig. 4 and conductivity in Fig. 3. The peak at 406 cm − 1 may be characteristic of some oxygen rich sulfide crystal or unreacted Li2O, as its presence is shown in the results of X-ray diffraction. The –30P2S5 system predominantly shows formation of the unknown crystal for high values of x, but exhibits formation of the Li4P2S6 crystal at x below 35, at the wavenumber 387 cm − 1, with some formation of the Li3PS4 crystal for increasing Li2O content until x = 10. As the Li4P2S6 crystal is not as highly conducting as the Li3PS4 crystal, this result is consistent with the observation of reduced ionic conductivity in Fig. 3 [45]. These Raman measurements confirm the formation of Li3PS4 and Li4P2S6 crystals as well as a new crystal in the electrolyte, and show the dependence of crystal formation on composition of electrolyte. As mentioned previously, sulfide based solid electrolytes (thio-LISICONs) are preferable over oxide based solid electrolytes (LISICONs) for their high polarizability and large ionic radius that allow for higher conductivity. However, sulfide based electrolytes are not completely stable against oxide based cathode materials such as LiCoO2 and typically result in rapid capacity fade. This degradation is attributed to the large potential difference between the sulfide and oxide based species. Previously, it was proposed that the degradation is a result of the formation of a space charge layer in which a resistance layer develops as a result of mixed ion conduction in the electrode material and li-ion conduction in the electrolyte resulting in depletion of the lithium in the electrolyte [38,46]. Here, we propose that because sulfur is much less electronegative than oxygen, it is more likely to form compounds in which it has a positive oxidation number such as SO2 or even SO3 as shown by Eqs. (1)–(5), which can then bind to a transition metal. Sulfur dioxide for instance, can bind to metal ions as a ligand to form a metal S þ O2 →SO2
ð1Þ
Fig. 5. Raman spectroscopic measurements of the –20P2S5, –25P2S5, and –30P2S5 solid electrolyte systems for increasing value of x by increments of 5.
S þ 3=2O2 →SO3 ΔH ¼ −395:8kJ=mole
ð2Þ
2SO2 þ O2 →2SO3 ΔH ¼ −198:2 kJ=mole
ð3Þ
SO3 →SO2 þ 1=2O2 ΔH ¼ þ99:1 kJ=mole
ð4Þ
S þ O2 →SO2 ΔH ¼ −296:7 kJ=mole
ð5Þ
sulfur dioxide complex, typically with transition metals in a 0 or + 1 oxidation state [47,48]. In regards to LiCoO2, the transition metal Co contained in the active material is consumed in such a reaction by the electrolyte upon delithiation and converted to an inactive sulfur containing compound such as CoS. Uchimoto et al. have previously indicated the presence of CoS through X-ray absorption spectroscopy at the interface of the LiCoO2 electrode and Li2S–P2S5 solid electrolyte
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[49]. Degradation of the capacity may also be a result of active material diffusion into the electrolyte, proposed by Sakuda et al. in which an interfacial layer comprised of Co and S develops during charging and discharging [50]. Degradation may be a result of one or all three of either space-charge resistance, diffusion of active materials, or consumption of active materials. Fig. 6 shows the charge–discharge voltage profiles for LiCoO2/Li cells using numerous different compositions of electrolyte to show the variations in first cycle side reaction. A strong correlation of reduced first cycle side reaction with increased Li2O content is observed for all 3 data sets. The –20P2S5 system shows the best overall performance with very little overall side reaction and a reversible capacity of almost 90 mAh g− 1 for the sample x = 45. The –25P2S5 system showed little change in electrochemical profile for the initial charge, but exhibited much better reversible capacity for higher concentrations
of Li2O in the electrolyte. The –30P2S5 system showed very poor performance which is attributed to lower conductivity and mixed crystallites in the electrolyte that react irreversibly with the electrode material. Solid electrolytes containing Li2O exhibit greatly reduced first cycle side reactions that result in higher reversible capacity. Fig. 7 shows the cycling behavior of LiCoO2/Li cells utilizing the same electrolytes displayed in Fig. 6 with a trend of higher reversible capacity and better cycle stability with increasing Li2O content. Comparison to our previous work exhibiting performance of the binary 80Li2S–20P2S5 solid electrolyte shows a significant improvement of the –20P2S5 series with added Li2O in this work [36]. The enhanced cycle stability and higher capacity is attributed to the formation of highly conducting, and Li stable electrolyte compounds as shown in Figs. 4 and 5. As previously shown high concentrations of Li2O increase resistance and therefore result in a nominal capacity, however in the exceptional case of x = 20 of the –30P2S5 system with the
Fig. 6. First cycle voltage profiles for x = 45, 55, and 65 of the –20P2S5; x = 65, 70, and 75 of the –25P2S5; and x = 60, 65, and 70 of the –30P2S5.
Fig. 7. Cycle life behavior for x = 45, 55, and 65 of the –20P2S5; x = 65, 70, and 75 of the –25P2S5; and x = 60, 65, and 70 of the –30P2S5.
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heightened conductivity, no observable capacity was obtained. We attribute this to the active material loss by one of the previously explained mechanisms. With a reversible capacity of only 50 mAh g − 1 after 25 cycles, even the –20P2S5 electrolytes do not exhibit a capacity high enough for commercialization. While these data are proof of concept, further research is necessary to fully stabilize the electrode/electrolyte interaction. 4. Conclusion Li2S–Li2O–P2S5 solid state electrolyte systems with varying composition were prepared by planetary ball milling. The xLi2S-(80-x) Li2O–20P2S5 performed best with the highest conductivity of tested electrolytes at 6.5 × 10 − 5 S cm − 1 and the best cycle performance achieving almost 90 mAh g − 1 reversible capacity on the second cycle at x = 45. Structural analysis showed that inclusion of Li2O into the Li2S–P2S5 system resulted in mixed crystalline species that lead to lower conductivity, but heightened reversible capacity and cycle stability. While addition of Li2O increased cyclability of cells with LiCoO2 cathodes, low overall capacities achieved by these electrolytes are not acceptable for rechargeable solid state batteries. Further study of the increased stability observed with the introduction of Li2O into the electrolyte may have future application for advanced solid state batteries. Verification of the newly proposed consumption reaction must be performed and investigation of the formation of potentially gaseous SO2 or SO3 in the electrolyte could help explain high resistance to lithium-ion migration in the sulfur based solid electrolytes. Currently studies with SSBM of the same Li2S–Li2O–P2S5 solid electrolytes are showing even greater reversible capacity and cycle stability. Acknowledgements The authors gratefully acknowledge Darpa DSO for funding this research. References [1] Y. Zhu, C. Wang, Department of Chemical and Biomolecular Engineering, University of Maryland, 2010, pp. 2830–2841. [2] T. Minami, A. Hayashi, M. Tatsumisago, Solid State Ionics 177 (2006) 2715. [3] M. Murayama, N. Sonoyama, A. Yamada, R. Kanno, Solid State Ionics 170 (2004) 173. [4] F. Mizuno, A. Hayashi, K. Tadanaga, M. Tatsumisago, Solid State Ionics 177 (2006). [5] A. Sakuda, A. Hayashi, T. Ohtomo, S. Hama, M. Tatsumisago, J. Power Sources 195 (2010) 599–603. [6] H. Kitaura, A. Hayashi, K. Tadanaga, M. Tatsumisago, J. Power Sources 189 (2009). [7] A. Robertson, A. West, A. Ritchie, Solid State Ionics 104 (1997). [8] H. Muramatsu, et al., Solid State Ionics (2010), doi:10.1016/j.ssi.2010.10.013.
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Solid State Ionics journal homepage: www.elsevier.com/locate/ssi
Li2S–Li2O–P2S5 solid electrolyte for all-solid-state lithium batteries James E. Trevey ⁎, Jeremy R. Gilsdorf, Sean W. Miller, Se-Hee Lee Department of Mechanical Engineering, University of Colorado-Boulder, Boulder, CO 80309-0427, USA
a r t i c l e
i n f o
Article history: Received 2 November 2011 Received in revised form 17 February 2012 Accepted 20 February 2012 Available online 28 March 2012 Keywords: Solid-state Lithium Battery Sulfur Lithium cobalt oxide
a b s t r a c t Various compositions of Li2S–Li2O–P2S5 solid electrolytes were prepared by planetary ball milling and systematically investigated. Structural analysis by X-ray diffraction (XRD) and Raman spectroscopy showed phase changes in the electrolyte corresponding to changes in conductivity. All-solid-state LiCoO2/Li cells were tested by constant–current–constant–voltage (CCCV) charge–discharge cycling at a current density of 50 μA cm− 2 between 2.5 and 4.3 V (vs. Li/Li+). In spite of reduced conductivity from introducing oxide species into a sulfide based electrolyte, increased cycle stability is observed for increased concentrations of Li2O in the electrolyte. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Advancements in energy storage technology for applications ranging from portable electronics to solar energy are continually sought by modern science. Lithium-ion batteries have always attracted considerable attention as rechargeable power sources due to high achievable energy and power densities, as well as good shelf life and no memory effect [1–7]. However, typical lithium-ion batteries containing liquidelectrolyte maintain numerous safety issues regarding flammability and toxicity [8–14]. These safety issues not only require considerable “dead weight” in regards to the devices and insulation required for controlling and containing the hazardous liquid which reduces energy density, but restrict scale up for potential applications regarding PHEV and EV due to the higher risk associated with increased size of liquidelectrolyte batteries [1,6,15–20]. Therefore improving safety of rechargeable lithium-ion batteries has become of great interest to the research community with respect to overall safety and performance [10,21]. This drives research for improvement of all-solid-state batteries which are inherently safer, as they do not contain any liquids, exhibit leakage, or have potential to undergo thermal runaway [5,8,12,14,18,20,22]. Though currently, all-solid-state batteries cannot match the performance of liquid-based batteries due to low achievable ionic conductivity and interfacial instability of solid electrolytes with electrode materials [4,10,13,23,24]. A large variety of solid electrolytes have been identified for use in all-solid-state batteries with ranging properties and target characteristics. The sulfide based electrolyte Li2S–P2S5 has been of continued ⁎ Corresponding author at: HRL Laboratories LLC, Malibu, CA 90265-4797, USA. Tel.: +1 310 317 5137; fax: +1 310 317 5840. E-mail address: [email protected] (J.E. Trevey). 0167-2738/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2012.02.034
research for over two decades for its high ionic conductivity, stability in contact with lithium metal, and simplicity in design [3,11,15,17,24–27]. Numerous additives to the electrolyte such as GeS2, GeSe2, P2O5, P2S3, SiS2 and more have been introduced to the electrolyte to aid in interfacial stabilization of solid electrolytes with positive electrode materials such as LiCoO2 [2,6,23,28]. However, research has shown limited success in stabilizing the capacity fade of solid batteries with sulfide based electrolytes and oxide based positive electrode materials. The difference in potential between these species causes the development of highly resistive reaction products within batteries that results in rapid capacity fade [5,10–12,14,20–22]. Opposed to melt and quench techniques that have been employed in the past; the use of ball milling as a means of producing solid electrolytes offers numerous advantages [11,25]. From the standpoint of manufacturability, ball milling is lower cost and less time consuming, while from a materials science aspect, ball milling is capable of producing ultra-fine powders sufficient for high conductivity of solid electrolyte and good interfacial contact between electrolyte and electrode materials [14,20]. Ball milling embraces a complex mixture of fracturing, grinding, high-speed plastic deformation, cold welding, thermal shock, and intimate mixing [26–30]. The mechanical energy created by the milling can induce chemical reactions of the reactants at room temperature [31], introduce a large number of crystal defects, and increase the specific surface area of compounds, thought to be one of the main reasons for increased reactivity of milled materials [32,33]. The present paper reports on the ternary solid electrolyte system Li2S–Li2O–P2S5 generated by planetary ball milling regarding characterization and application to all-solid-state secondary batteries. In an effort to reduce the potential between the sulfide based electrolyte and oxide based cathode, LiCoO2, an oxygen containing species, Li2O, was introduced to the electrolyte. In concept, use of the ternary
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electrolyte having sulfide and oxide lithium-containing species will result in reduced side reactions and improved cycling stability of batteries. In order to get a global view of the system, three sets of data with varying compositions were manufactured, characterized, and tested. With the typical range of x = 70–80 in the binary xLi2S-(100x)P2S5 system for highly conducting and stable electrolyte, we used a similar compositional scheme in which the molar amount of P2S5 was fixed, and the lithium containing compound ratios were varied. Our results detail an investigation of the ternary electrolyte in the ranges of xLi2S-(80-x)Li2O–20P2S5 for x = 0–80 (−20P2S5), xLi2S-(75-x)Li2O–25P2S5 for x = 0–75 (−25P2S5), and xLi2S-(70-x) Li2O–30P2S5 for x = 0–70 (−30P2S5). 2. Experimenatal Li2S–Li2O–P2S5 solid electrolytes (SE) were prepared by planetary ball milling. Reagent Grade Powders of Li2S (Alderich, 99.999%), P2S5 (Alderich 99%) and Li2O (Alfa-Aesar 99.9%) were used as starting materials. Appropriate concentrations of materials were combined into a stainless steel agate vial. SE was prepared using 2 large stainless steel balls (10 mm diameter) and 6 small stainless steel balls (6 mm diameter) for grinding. High energy ball milling (Across International PQ-N4) took place for 20 continuous hours in a sealed Ar container. Composite electrodes were prepared by mixing LiCoO2 powder (Sigma Alderich), SE Li2S–Li2O–P2S5 and acetylene black (Alfa-Aesar, 50% compressed) at a weight ratio of 20:30:3 respectively. Bilayer electrolyte pellets are formed by hand pressing 100 mg of Li2S–Li2O– P2S5 on top of a 100 mg hand pressed layer of 77.5Li2S-22.5P2S5 (mol%) SE. A 10 mg layer of the composite material is then carefully spread on the top of the Li2S–Li2O–P2S5 electrolyte layer and the cell pelletized by cold pressing (5 metric tons) for 5 min. Li foil (AlfaAesar, 0.75 mm thick) is then attached to the 77.5Li2-22.5P2S5 (mol%) SE face that was previously shown to be stable against Li metal [34], at 2 metric tons. All pressing and testing operations are carried out in a polyarytherketone (PEEK) mold (ϕ = 1.3 cm) with Ti metal rods as current collectors for both working and counter electrodes. All processes were carried out in an Ar-filled glove box. Galvanostatic charge–discharge cycling took place at first cycle cut off voltages of 4.3 and 2.5 V at a current of 50 μA cm − 2 at room temperature using an Arbin BT2000. SE samples were characterized by XRD measurement with Cu-Kα radiation. Samples were sealed in an airtight aluminum container with Beryillium windows and mounted on the X-ray diffractometer (PANalytical, PW3830). Raman spectroscopic measurments were taken through sealed glass containers (Jasco NRS-3100). Ionic conductivities were measured by AC impedance spectroscopy (Solarton 1280 C) of which weighed materials are cold pressed to both sides
of the SE pellet to serve as electrodes. The impedance of selected cells was measured from 20 MHz to 100 mHz at room temperature and the conductivity was determined using complex impedance analysis. Schematic diagrams for the Li/SE/LiCoO2 cells and AC impedance cells can be seen Fig. 1(a) and (b), respectively. 3. Results Although sulfide based solid electrolyte (thio-LISICONs) are preferable over oxide based solid electrolytes (LISICONs) for their high polarizability and large ionic radius that allow for higher conductivity, sulfide based electrolytes are not completely stable against oxide based cathode materials such as LiCoO2 due to the large potential difference. While the introduction of an oxide-based species, Li2O, into the sulfide based solid electrolyte is expected to cause a loss of conductivity, it is proposed here that incorporation of Li2O into the binary Li-P-S electrolyte will mitigate the potential difference between the oxide-based cathode and sulfide based electrolyte, thereby reducing any side reactions that may result in increased internal resistance within the cell. Fig. 2 shows impedance data for two cells made with different solid electrolytes, one rich in Li2O and one poor in Li2O. The data depict the impedance measurements immediately after cell fabrication, after first charge, and after first discharge. As expected, higher overall impedance is observed for the sample with more Li2O. However, the normalized difference from the chargeinterfacial impedance and the discharge-interfacial impedance of the two samples shows a 6% reduction of interfacial impedance for the Li2O rich cell. Increased stability of the electrode/electrolyte contact is projected to outweigh the conductivity losses of the solid electrolyte. Fig. 3 shows a comparison of the conductivities of the three ball milled electrolyte systems. A generally decreasing conductivity can be observed for decreasing x, or decreasing Li2S content and increasing Li2O content, as expected. Furthermore, with an increased molecular composition of P2S5, a similar trend of decreasing conductivity can be observed, for instance, with the maximum conductivity of the –20P2S5, –25P2S5, and the –30P2S5 systems being 3.0 × 10− 4 S cm- 1, 1.6 × 10− 4 S cm- 1, and 5.4 × 10− 5 S cm − 1 respectively, resulting from smaller concentrations of Li mobile ions in the electrolyte [35,36]. This trend agrees with our previous work regarding single step ball milling (SSBM), but with lower overall conductivity [36]. Machida et al. also tested the Li2S–Li2O–P2S5 system and found a high conductivity of 2.7 × 10− 4 S cm− 1 for a composition of 67.5Li2S–7.5Li2O–25P2S5, of which composition we calculate a conductivity of 1.1 × 10− 4 S cm− 1 for our electrolyte [37]. The difference in conductivity values is relatively small and likely due to minute differences in milling conditions. While the decreasing conductivity of the –20P2S5 system with decreasing x remains relatively consistent, the –25P2S5 system shows a region
Fig. 1. Schematic diagram of Ti test die for a) the all-solid-state battery and b) measuring conductivity of solid-electrolyte.
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27
are used for standardization of the data. Throughout all 3 groups of electrolyte, highly crystalline peaks for Li2S at 26.8° are observed at high values of x, while peaks at 34° for Li2O are observed for lower values of x [38,39]. A broad amorphous peak is observed for all 3 sample groups from 25° to 35°, with the –20P2S5 samples showing a second broad peak at 17°–18° that briefly appears for high values of x in the – 25P2S5 patterns as well. While planetary ball milling typically results in amorphous electrolyte, our electrolyte exhibits some crystalline behavior as a result of the increased temperature of the ball milling vials during the 20 consecutive hours of rotation, similar to the SSBM formation of highly crystalline and highly conducting materials by ball milling in a high temperature environment [34].
Fig. 2. Impedance for (a) x = 50 and (b) x = 75 of the xLi2S-(80-x)Li2O–20P2S5 electrolyte system.
of markedly lower conductivity between x = 35–50, and the –30P2S5 system shows a rapid decrease in conductivity as x approaches zero with the exception of one point at x = 20. For both the –20P2S5 and the –25P2S5 systems there is approximately 1 order of magnitude difference in conductivity between x = maximum, and x = 0 while the –30P2S5 system shows a decrease in conductivity of 6 orders of magnitude. X-ray diffraction patterns for the electrolytes can be observed in Fig. 4 with significant differences between the three systems. Shown by group according to the molar amount of P2S5, a distinct decrease in relative intensity is observed for increasing P2S5 content. Small peaks at 25° and 38° are consistent with records for Be and
Fig. 3. Conductivity map of the –20P2S5, –25P2S5, and –30P2S5 solid electrolyte systems.
Fig. 4. X-ray diffraction comparison of the –20P2S5, –25P2S5, and –30P2S5 solid electrolyte systems for increasing value of x by increments of 5.
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Interestingly, comparison between the XRD patterns and the conductivity map from Fig. 3 exhibits significant correlation between data sets. In regards to the –25P2S5 system, a substantial intensity and crystallographic difference for the region x = 35–50 as was observed by a marked decrease in conductivity for the same range of x in Fig. 3. In this data range, the development of a broad peak around 16° which has been previously attributed to a thio-LISICON III analog [40] is observed, as well as a highly crystalline peak corresponding to the presence of Li2O. Based on the results of XRD, the drop in conductivity of the –25P2S5 system is a result of increased Li2O presence in the electrolyte. Similarly for the –30P2S5 system, at the composition x = 20, a sharp increase of conductivity was observed in Fig. 3, which correlates to the disappearance of the Li2O peak in the corresponding XRD pattern. The Raman spectra of Li3PS4 and Li4P2S6 crystals have been previously displayed by Machida et al. [41]. The Li3PS4 spectrum exhibits a strong peak at 424 cm − 1 and two weak peaks at 550 and 567 cm − 1 of which Tachez et al. [42] ascribed to tetrahedral PS43- unit existing in the Li3PS4 crystal. The Li4P2S6 spectrum showed a strong peak at 387 cm − 1 and weak peaks at 550, 570, and 257 cm− 1 of which Mercier et al. [43,44] reported that such peaks can be attributed to PS3 terminal modes of P2S64 − unit; i.e., the P–P bonded hexathiohypophosphate group. In Fig. 5, Raman spectroscopic measurement of all three systems of materials reveals a number of different relationships between material compositions, X-ray diffraction patterns, and conductivity measurements. The –20P2S5 system predominantly showed evidence of the Li3PS4crystal, the intensity of which increases with increasing composition of Li2S, indicating that increased Li2S content leads to the formation of the PS43- unit and higher conductivity. The –25P2S5 system exhibited characteristics of the Li3PS4 crystal, as well an unknown peak near 406 cm - 1 at the midpoint between Li3PS4and Li4P2S6 crystals. A sharp shift from of the strong peak is observed near x = 35, similarly to X-ray diffraction in Fig. 4 and conductivity in Fig. 3. The peak at 406 cm − 1 may be characteristic of some oxygen rich sulfide crystal or unreacted Li2O, as its presence is shown in the results of X-ray diffraction. The –30P2S5 system predominantly shows formation of the unknown crystal for high values of x, but exhibits formation of the Li4P2S6 crystal at x below 35, at the wavenumber 387 cm − 1, with some formation of the Li3PS4 crystal for increasing Li2O content until x = 10. As the Li4P2S6 crystal is not as highly conducting as the Li3PS4 crystal, this result is consistent with the observation of reduced ionic conductivity in Fig. 3 [45]. These Raman measurements confirm the formation of Li3PS4 and Li4P2S6 crystals as well as a new crystal in the electrolyte, and show the dependence of crystal formation on composition of electrolyte. As mentioned previously, sulfide based solid electrolytes (thio-LISICONs) are preferable over oxide based solid electrolytes (LISICONs) for their high polarizability and large ionic radius that allow for higher conductivity. However, sulfide based electrolytes are not completely stable against oxide based cathode materials such as LiCoO2 and typically result in rapid capacity fade. This degradation is attributed to the large potential difference between the sulfide and oxide based species. Previously, it was proposed that the degradation is a result of the formation of a space charge layer in which a resistance layer develops as a result of mixed ion conduction in the electrode material and li-ion conduction in the electrolyte resulting in depletion of the lithium in the electrolyte [38,46]. Here, we propose that because sulfur is much less electronegative than oxygen, it is more likely to form compounds in which it has a positive oxidation number such as SO2 or even SO3 as shown by Eqs. (1)–(5), which can then bind to a transition metal. Sulfur dioxide for instance, can bind to metal ions as a ligand to form a metal S þ O2 →SO2
ð1Þ
Fig. 5. Raman spectroscopic measurements of the –20P2S5, –25P2S5, and –30P2S5 solid electrolyte systems for increasing value of x by increments of 5.
S þ 3=2O2 →SO3 ΔH ¼ −395:8kJ=mole
ð2Þ
2SO2 þ O2 →2SO3 ΔH ¼ −198:2 kJ=mole
ð3Þ
SO3 →SO2 þ 1=2O2 ΔH ¼ þ99:1 kJ=mole
ð4Þ
S þ O2 →SO2 ΔH ¼ −296:7 kJ=mole
ð5Þ
sulfur dioxide complex, typically with transition metals in a 0 or + 1 oxidation state [47,48]. In regards to LiCoO2, the transition metal Co contained in the active material is consumed in such a reaction by the electrolyte upon delithiation and converted to an inactive sulfur containing compound such as CoS. Uchimoto et al. have previously indicated the presence of CoS through X-ray absorption spectroscopy at the interface of the LiCoO2 electrode and Li2S–P2S5 solid electrolyte
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[49]. Degradation of the capacity may also be a result of active material diffusion into the electrolyte, proposed by Sakuda et al. in which an interfacial layer comprised of Co and S develops during charging and discharging [50]. Degradation may be a result of one or all three of either space-charge resistance, diffusion of active materials, or consumption of active materials. Fig. 6 shows the charge–discharge voltage profiles for LiCoO2/Li cells using numerous different compositions of electrolyte to show the variations in first cycle side reaction. A strong correlation of reduced first cycle side reaction with increased Li2O content is observed for all 3 data sets. The –20P2S5 system shows the best overall performance with very little overall side reaction and a reversible capacity of almost 90 mAh g− 1 for the sample x = 45. The –25P2S5 system showed little change in electrochemical profile for the initial charge, but exhibited much better reversible capacity for higher concentrations
of Li2O in the electrolyte. The –30P2S5 system showed very poor performance which is attributed to lower conductivity and mixed crystallites in the electrolyte that react irreversibly with the electrode material. Solid electrolytes containing Li2O exhibit greatly reduced first cycle side reactions that result in higher reversible capacity. Fig. 7 shows the cycling behavior of LiCoO2/Li cells utilizing the same electrolytes displayed in Fig. 6 with a trend of higher reversible capacity and better cycle stability with increasing Li2O content. Comparison to our previous work exhibiting performance of the binary 80Li2S–20P2S5 solid electrolyte shows a significant improvement of the –20P2S5 series with added Li2O in this work [36]. The enhanced cycle stability and higher capacity is attributed to the formation of highly conducting, and Li stable electrolyte compounds as shown in Figs. 4 and 5. As previously shown high concentrations of Li2O increase resistance and therefore result in a nominal capacity, however in the exceptional case of x = 20 of the –30P2S5 system with the
Fig. 6. First cycle voltage profiles for x = 45, 55, and 65 of the –20P2S5; x = 65, 70, and 75 of the –25P2S5; and x = 60, 65, and 70 of the –30P2S5.
Fig. 7. Cycle life behavior for x = 45, 55, and 65 of the –20P2S5; x = 65, 70, and 75 of the –25P2S5; and x = 60, 65, and 70 of the –30P2S5.
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heightened conductivity, no observable capacity was obtained. We attribute this to the active material loss by one of the previously explained mechanisms. With a reversible capacity of only 50 mAh g − 1 after 25 cycles, even the –20P2S5 electrolytes do not exhibit a capacity high enough for commercialization. While these data are proof of concept, further research is necessary to fully stabilize the electrode/electrolyte interaction. 4. Conclusion Li2S–Li2O–P2S5 solid state electrolyte systems with varying composition were prepared by planetary ball milling. The xLi2S-(80-x) Li2O–20P2S5 performed best with the highest conductivity of tested electrolytes at 6.5 × 10 − 5 S cm − 1 and the best cycle performance achieving almost 90 mAh g − 1 reversible capacity on the second cycle at x = 45. Structural analysis showed that inclusion of Li2O into the Li2S–P2S5 system resulted in mixed crystalline species that lead to lower conductivity, but heightened reversible capacity and cycle stability. While addition of Li2O increased cyclability of cells with LiCoO2 cathodes, low overall capacities achieved by these electrolytes are not acceptable for rechargeable solid state batteries. Further study of the increased stability observed with the introduction of Li2O into the electrolyte may have future application for advanced solid state batteries. Verification of the newly proposed consumption reaction must be performed and investigation of the formation of potentially gaseous SO2 or SO3 in the electrolyte could help explain high resistance to lithium-ion migration in the sulfur based solid electrolytes. Currently studies with SSBM of the same Li2S–Li2O–P2S5 solid electrolytes are showing even greater reversible capacity and cycle stability. Acknowledgements The authors gratefully acknowledge Darpa DSO for funding this research. References [1] Y. Zhu, C. Wang, Department of Chemical and Biomolecular Engineering, University of Maryland, 2010, pp. 2830–2841. [2] T. Minami, A. Hayashi, M. Tatsumisago, Solid State Ionics 177 (2006) 2715. [3] M. Murayama, N. Sonoyama, A. Yamada, R. Kanno, Solid State Ionics 170 (2004) 173. [4] F. Mizuno, A. Hayashi, K. Tadanaga, M. Tatsumisago, Solid State Ionics 177 (2006). [5] A. Sakuda, A. Hayashi, T. Ohtomo, S. Hama, M. Tatsumisago, J. Power Sources 195 (2010) 599–603. [6] H. Kitaura, A. Hayashi, K. Tadanaga, M. Tatsumisago, J. Power Sources 189 (2009). [7] A. Robertson, A. West, A. Ritchie, Solid State Ionics 104 (1997). [8] H. Muramatsu, et al., Solid State Ionics (2010), doi:10.1016/j.ssi.2010.10.013.
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