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Interfacial modification for high-power solid-state lithium batteries
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Solid State Ionics 179 (2008) 1333 – 1337 www.elsevier.com/locate/ssi
Interfacial modification for high-power solid-state lithium batteries Kazunori Takada⁎, Narumi Ohta, Lianqi Zhang, Katsutoshi Fukuda, Isao Sakaguchi, Renzhi Ma, Minoru Osada, Takayoshi Sasaki National Institute for Materials Science, 1-1, Namiki, Tsukuba, Ibaraki 305-0044, Japan CREST, Japan Science and Technology Agency, Japan Received 12 July 2007; received in revised form 4 February 2008; accepted 12 February 2008
Abstract Interfaces between LiCoO2 and sulfide solid electrolytes were modified in order to enhance the high-rate capability of solid-state lithium batteries. Thin films of oxide solid electrolytes, Li4Ti5O12, LiNbO3, and LiTaO3, were interposed at the interfaces as buffer layers. Changes in the high-rate performance upon heat treatment revealed that the buffer layer should be formed at low temperature to avoid thermal diffusion of the elements. Buffer layers of LiNbO3 and LiTaO3 can be formed at low temperature for the interfacial modification, because they show high ionic conduction in their amorphous states, and so are more effective than Li4Ti5O12 for high-power densities. © 2008 Elsevier B.V. All rights reserved. Keywords: Solid electrolyte; Lithium battery; Nanoionics; Space-charge layer; LiTaO3
1. Introduction Solid-state lithium batteries with inorganic solid electrolytes would eliminate the safety problem of lithium batteries arising from their combustible organic electrolytes. However, the use of a solid electrolyte reduces the power density, mainly because of poor ionic conduction in the solid electrolytes, so studies on solid-state lithium batteries have focused on enhancing their ionic conductivities. Recently, highly ion-conductive sulfides have been developed [1,2] with ionic conductivities and activation energies for conduction of the order of 10− 3 S cm− 1 and below 0.2 eV, respectively. However, even when such highly-conductive solid electrolytes were used, the high-rate capabilities of solid-state lithium batteries had not been remarkably improved [3,4]. Because ionic conduction in solid electrolytes is as fast as that in liquid electrolytes, and electrodes used in such studies were LiCoO2 and graphite, which are the
same as those found in lithium–ion cells, it is conjectured that the rate-determining step is not in the bulk of the materials, but at the interface between the electrode and the electrolyte. We therefore focused on the interface and succeeded in enhancing the high-rate capability of solid-state lithium batteries by modification of the interface. Thin buffer layers of Li4Ti5O12 [5] or LiNbO3 [6] were interposed between LiCoO2 and the sulfide solid electrolytes, which drastically reduced the interfacial resistance and increased the current drain. The improvement of the high-rate capability was further in the case of LiNbO3, which was attributed to its higher ionic conductivity. In this study, LiTaO3 was used as another buffer layer material, and the electrode performance was compared with those of Li4Ti5O12 and LiNbO3 in order to confirm it, because ionic conductivity of LiTaO3 is higher than Li4Ti5O12 and as high as LiNbO3 [7]. In additions, some details in our previous studies are presented. 2. Experimental
⁎ Corresponding author. National Institute for Materials Science, 1-1, Namiki, Tsukuba, Ibaraki 305-0044, Japan. E-mail address: [email protected] (K. Takada). 0167-2738/$ - see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2008.02.017
In our previous studies, the buffer layers were formed on the surface of LiCoO2 particles in order to interpose them between
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K. Takada et al. / Solid State Ionics 179 (2008) 1333–1337
the LiCoO2 and the sulfide electrolyte, when they were mixed into cathodes. Similarly, in the present study, the LiCoO2 particles were coated with LiTaO3 by spray coating as described below. First, a coating solution of LiTaO3 was prepared. Li metal (1.47 g, 0.21 mol) was dissolved in anhydrous ethanol to obtain an ethanol solution of lithium ethoxide. The solution was mixed with tantalum pentaethoxide (86.26 g, 0.21 mol) and diluted by anhydrous ethanol to make the weight of the final solution 700 g. The solution was sprayed onto the surface of LiCoO2 powder (D10, Toda Kogyo) by a rolling fluidized coating machine (MP-01, Powrex). Different amounts of the solution were sprayed in order to control the thickness of the buffer layer. After the coating, the samples were heated at 400 °C for 30 min under an oxygen flow to decompose the organic components. The samples were labeled by the thickness of the buffer layer, t, in this paper, which was calculated from the amounts of the applied solution and the BET surface area of the LiCoO2 powder (0.26 m2 g− 1). The thickness of the buffer layer was confirmed by transmission electron microscopy. The electrode properties of the samples were investigated in two-electrode cells by the same procedure as described in our previous papers [5,6]. In–Li alloy was used as the counter electrode in the electrochemical measurements, which has an electrode potential of 0.62 V vs. Li+/Li. For clarity, the cell voltages are presented by adding 0.62 V as if lithium metal was used as the counter electrode. Li3.25Ge0.25P0.75S4 was used as the solid electrolyte. Since it is difficult to characterize thin buffer layers on the surface of the LiCoO2 particles, thin films prepared by spin coating and pulse laser deposition (PLD) were characterized instead. X-ray diffraction (XRD) patterns for the spin-coated films were taken in order to investigate the crystal structures of the buffer layers. The ethanol solution of the alkoxides mentioned above and in Refs. [5] and [6] were spin-coated onto Si (100) substrates, and the coated films were then heated at different temperatures for 30 min under oxygen atmosphere. In-plane XRD data were collected for the thin films using synchrotron radiation at Photon Factory BL-3A at the Institute of Materials Structure Science, High Energy Accelerator Research Organization (KEKPF). The wavelength of the incident X-ray was calibrated to be 1.39904(8) Å. An NaI scintillation counter was scanned to take the diffraction data with a grazing exit angle of 0.2° around the normal axis of the thin films. The effect of heat treatment on the buffer layers was investigated using LiCoO2/Li4Ti5O12-stacked films prepared by PLD. LiCoO2 and Li4Ti5O12 powders synthesized by solid-state reactions were sintered into pellets, which were used as targets in the PLD. First, an LiCoO2 film was deposited onto an Si substrate heated at 600 °C using a KrF excimer laser (λ = 248 nm) operated at 10 Hz. Then, an Li4Ti5O12 film was deposited onto the LiCoO2 without heating the substrate to obtain LiCoO2/Li4Ti5O12-stacked films. The stacked films were heated at temperatures ranging from 400 °C to 600 °C for 10 min. Changes in the buffer layer upon the heat treatment were investigated by secondary ion mass spectrometry (SIMS; IMS 4f, CAMACA). Depth profiles for the films heated at different temperatures were measured.
3. Results and discussion 3.1. Strategy for the enhancement of high-rate capability When LiCoO2 and graphite are used as the cathode and the anode, respectively, the rate-determining step is the cathode. As shown in Fig. 1, discharge capacity decreased quickly for LiCoO2 with current density increased, whereas such drastic decrease was not observed for the graphite. In our previous study [5], the reason for the poor high-rate capability of the cathode was considered to be a highly-resistive space-charge layer formed at the interface between the LiCoO2 and the sulfide electrolyte. Electrochemical potentials of the lithium ions should be largely different between oxides and sulfides [8], because oxide ions attract lithium ions much more strongly than sulfide ions. This large difference will cause some of the lithium ions in the sulfide electrolyte to transfer to the LiCoO2 until the equilibrium reached, when they are brought into contact with each other, leaving a lithium-deficient space-charge layer on the electrolyte side at the interface. Since solid electrolytes are optimized in the compositions for fast ionic conduction, the changing lithium content will decrease the ionic conductivity in the space-charge layer. However, contact between oxides and sulfides alone does not always form highly-resistive layer [9]. An even greater problem is electronic conduction in LiCoO2. Interfacial phenomena at the hetero-junction of ionic conductors are now categorized into “nanoionics” [10], in which space-charge layer models for the hetero-junction between ionic conductors are analogous to those at a semiconductor interface. The space-charge layers at the LiCoO2/electrolyte interface will be largely developed as in the Schottky junction because of the electronic conduction in LiCoO2 [11], which will greatly suppress the ionic conduction at the interface. Therefore,
Fig. 1. Discharge curves for LiCoO2 in thio-LISICON (a) [5] and graphite electrodes in a Li2S–P2S5 glass-ceramics electrolyte (b) [4].
K. Takada et al. / Solid State Ionics 179 (2008) 1333–1337
we interposed thin buffer layers of oxides with ionic conduction and electronic insulation, i.e. oxide solid electrolytes, between the LiCoO2 and the solid electrolyte. The interposition forms two interfaces: one between the LiCoO2 and the oxide solid electrolyte, and the other between the oxide and the sulfide electrolyte. The space-charge layers will not be largely developed at both interfaces, because the former interface is between two oxides, between which the electrochemical potential of lithium ions will not be largely different. The latter consists of electronic insulators, where the space-charge layer will be less developed as at p–n junctions between semiconductors [11]. Therefore, we used Li4Ti5O12 as the buffer layer material in our first study [5]. Evidence for the lithium–ion transfer causing large development of the space-charge layer was found in the voltage
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profile at the very beginning of the first charging. Fig. 2 shows change in the voltage profile of the LiCoO2 by the interposition of Li4Ti5O12. When the LiCoO2 was not coated with Li4Ti5O12, the charge curve showed a slope prior to the 4 V plateau. Because some lithium ions should have been transferred from the solid electrolyte to the LiCoO2 during the formation of the space-charge layer, the LiCoO2 should accommodate excess lithium ions. They are considered to be deintercalated in the beginning of the first charging, giving the additional oxidative deintercalation steps appearing as the slopes. The slopes gradually became shorter with increasing amount of applied Li4Ti5O12, which indicates that the increasing coverage of the buffer layer gradually decreased the area where the space-charge layer was developed. Correspondingly, the interfacial resistance appearing as an arc decreased and reached
Fig. 2. Voltage profiles at the beginning of the first charging (left) and complex impedance plots for the cathodes. Dotted, dashed, and solid lines are for Li4Ti5O12, LiNbO3, and LiTaO3 coated LiCoO2, respectively. The complex impedance plots were obtained by subtracting the impedance of the electrolyte layers.
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a minimum at t = 5 nm, although the slope still remained. This means that the buffer layer was so thick that the resistance of the buffer layer itself became predominantly large, when a sufficient amount of Li4Ti5O12 to completely cover the surface was applied. This fact suggests that more conductive buffer layers should be used to improve the high-rate capability. In addition, another requirement for the buffer layers is that they should be formed by low-temperature heat treatment as described below. 3.2. Heat treatment for the formation of the buffer layer Examples of such highly-conductive oxides are Li1 + xMxTi1 − x (PO4)3 with NASICON structure [12] and (Li, La)TiO3 with perovskite structure [13]. However, heat treatment at high temperatures is necessary to attain high ionic conductivity. For instance, LiTi2(PO4)3 thin films prepared by a sol–gel method showed fast ionic conduction only in the crystalline phase after heating above 600 °C [14]. However, buffer layers formed at such high temperatures did not improve the high-rate capability. Fig. 3 compares charge–discharge curves for the LiCoO2 with and without the buffer layer of Li4Ti5O12. When the sample was heated at 600 °C, the buffer layer did not cause any changes in the electrode properties. On the other hand, when the heat treatment temperature was 400 °C, the electrode performance was improved. Changes in the buffer layer upon the heat treatment were investigated by in-plane XRD and SIMS. Fig. 4a shows inplane XRD patterns for the Li4Ti5O12 films prepared by spin coating heated at different temperatures. The film heated at 400 °C did not give any sharp reflections, which suggested the amorphous nature of the film. When the film was heated at 600 °C, sharp reflections appeared, indicating crystallization of the film. That is, the buffer layer became ineffective by the heat treatment at 600 °C, although it was crystallized.
Fig. 3. Charge–discharge curves for uncoated (upper) and Li4Ti5O12-coated LiCoO2 heated at 400 °C and 600 °C at a current density of 127 μA cm− 2.
Fig. 4. Synchrotron radiation in-plane diffraction patterns for the thin films prepared on Si (100) substrates by spin-coating. (a) Patterns for Li4Ti5O12 heated at different temperatures. The reflections were indexed on a cubic lattice with a lattice constant of a = 8.311(5) Å. (b) Pattern for LiTaO3 prepared at 400 °C. The reflections were indexed on a rhombohedral lattice with lattice constants of a = 5.142(2) Å and c = 13.87(1) Å.
The reason can be seen in the depth profiles obtained for the LiCoO2/Li4Ti5O12-stacked films shown in Fig. 5. These depth profiles clearly indicated that the heat treatment at high temperature caused cobalt ions to diffuse into the Li4Ti5O12 layer, even when the heating duration was only 10 min and the thickness of the Li4Ti5O12 layer was ca. 150 nm. The cobalt ions penetrating the Li4Ti5O12 layer should induce electronic conduction in it, and thus the buffer layer can not adequately suppress the development of the space-charge layer. Therefore, we used LiNbO3 in place of Li4Ti5O12 to improve the high-rate capability [6]. Because LiNbO3 has high ionic conduction in the
Fig. 5. Depth profiles of LiCoO2/Li4Ti5O12-stacked films heated at different temperatures. The vertical axis indicates the intensity ratio of Co/6Li measured by SIMS.
K. Takada et al. / Solid State Ionics 179 (2008) 1333–1337
amorphous state [7], the heating temperature for the buffer layer formation can be reduced. LiTaO3 was also reported to have fast ionic conduction in the amorphous state as LiNbO3, and thus LiTaO3 was used as the buffer layer in the present study. A thin film of LiTaO3 prepared by spin coating and heated at 400 °C gave reflections in the inplane XRD pattern as shown in Fig. 4a. However, halos were also observed, indicating that at least part of the film remained amorphous. 3.3. Improvement of high-rate capability by the interposition of lithium tantalum oxide buffer layer Electrode properties of the LiTaO3-coated LiCoO2 are presented with those of the Li4Ti5O12 and LiNbO3-coated one in Fig. 2. The slopes, which correspond to the capacities of the space-charge layers, were in good agreement with those observed for Li4Ti5O12 and LiNbO3. That is, the change depended only on the applied amount of buffer layer materials, and not on the kind of materials. If the increasing thickness is the reason for the gradual change, different buffer materials should cause different changes. Therefore, it is concluded again that the increasing coverage is responsible for the gradual shortening of the slope. On the other hand, the LiTaO3 reduced the interfacial resistance more prominently than the Li4Ti5O12 and equivalently to the LiNbO3; the minimum interfacial resistance was almost the same as that observed for the LiNbO3 system. 4. Conclusions As presented in this paper, oxide solid electrolytes introduced to solid-state batteries with sulfide electrolytes as buffer layers improved the high-rate capability; that is, two kinds of solid electrolytes were used to enhance the performance. In our previous study on improving energy density of solid-state lithium batteries [3], we used two kinds of solid electrolytes in a battery: one resistive to highly-oxidizing
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LiCoO2 and the other resistive to highly-reducing C6Lix. These studies may be categorized as a “multi-electrolyte system”, which is valid only in solid-state batteries and effective for improving battery performance. Acknowledgements The authors thank T. Kurihara of Powrex Corp. for his help in preparing the coated LiCoO2. This work was partially supported by the Ministry of Economy, Trade, and Industry (MITI), New Energy, Industrial Technology Development Organization (NEDO), and the Grant-in-Aid for Scientific Research on Priority Area, “Nanoionics (439)” (19017018) by the Ministry of Education, Culture, Sports, Science and Technology (MEXT) of Japan. References [1] R. Kanno, M. Murayama, J. Electrochem. Soc. 148 (2001) A742. [2] F. Mizuno, A. Hayashi, K. Tadanaga, M. Tatsumisago, Adv. Mater. 17 (2005) 918. [3] K. Takada, T. Inada, A. Kajiyama, H. Sasaki, S. Kondo, M. Watanabe, M. Murayama, R. Kanno, Solid State Ionics 158 (2003) 269. [4] Y. Seino, K. Takada, B.-C. Kim, L. Zhang, N. Ohta, H. Wada, M. Osada, T. Sasaki, Solid State Ionics 176 (2005) 2389. [5] N. Ohta, K. Takada, L. Zhang, R. Ma, M. Osada, T. Sasaki, Adv. Mater. 18 (2006) 2226. [6] N. Ohta, K. Takada, I. Sakaguchi, L. Zhang, R. Ma, K. Fukuda, M. Osada, T. Sasaki, Electrochem. Commun. 9 (2007) 1486. [7] A.M. Glass, K. Nassau, T.J. Negran, J. Appl. Phys. 49 (1978) 4808. [8] P.G. Bruce, Chem. Commun. (1997) 1817. [9] K. Takada, M. Tansho, I. Yanase, T. Inada, A. Kajiyama, M. Kouguchi, S. Kondo, M. Watanabe, Solid State Ionics 139 (2001) 241. [10] J. Maier, Nat. Matters 4 (2005) 805. [11] J. Maier, Prog. Solid State Chem. 23 (1995) 171. [12] H. Aono, E. Sugimoto, Y. Sadaoka, N. Imanaka, G. Adachi, J. Electrochem. Soc. 136 (1989) 590. [13] Y. Inaguma, L. Chen, M. Itoh, T. Nakamura, T. Uchida, M. Ikuta, M. Wakihara, Solid State Commun. 86 (1993) 689. [14] K. Takada, K. Fujimoto, T. Inada, A. Kajiyama, M. Kouguchi, S. Kondo, M. Watanabe, Appl. Surf. Sci. 189 (2002) 300.
Solid State Ionics 179 (2008) 1333 – 1337 www.elsevier.com/locate/ssi
Interfacial modification for high-power solid-state lithium batteries Kazunori Takada⁎, Narumi Ohta, Lianqi Zhang, Katsutoshi Fukuda, Isao Sakaguchi, Renzhi Ma, Minoru Osada, Takayoshi Sasaki National Institute for Materials Science, 1-1, Namiki, Tsukuba, Ibaraki 305-0044, Japan CREST, Japan Science and Technology Agency, Japan Received 12 July 2007; received in revised form 4 February 2008; accepted 12 February 2008
Abstract Interfaces between LiCoO2 and sulfide solid electrolytes were modified in order to enhance the high-rate capability of solid-state lithium batteries. Thin films of oxide solid electrolytes, Li4Ti5O12, LiNbO3, and LiTaO3, were interposed at the interfaces as buffer layers. Changes in the high-rate performance upon heat treatment revealed that the buffer layer should be formed at low temperature to avoid thermal diffusion of the elements. Buffer layers of LiNbO3 and LiTaO3 can be formed at low temperature for the interfacial modification, because they show high ionic conduction in their amorphous states, and so are more effective than Li4Ti5O12 for high-power densities. © 2008 Elsevier B.V. All rights reserved. Keywords: Solid electrolyte; Lithium battery; Nanoionics; Space-charge layer; LiTaO3
1. Introduction Solid-state lithium batteries with inorganic solid electrolytes would eliminate the safety problem of lithium batteries arising from their combustible organic electrolytes. However, the use of a solid electrolyte reduces the power density, mainly because of poor ionic conduction in the solid electrolytes, so studies on solid-state lithium batteries have focused on enhancing their ionic conductivities. Recently, highly ion-conductive sulfides have been developed [1,2] with ionic conductivities and activation energies for conduction of the order of 10− 3 S cm− 1 and below 0.2 eV, respectively. However, even when such highly-conductive solid electrolytes were used, the high-rate capabilities of solid-state lithium batteries had not been remarkably improved [3,4]. Because ionic conduction in solid electrolytes is as fast as that in liquid electrolytes, and electrodes used in such studies were LiCoO2 and graphite, which are the
same as those found in lithium–ion cells, it is conjectured that the rate-determining step is not in the bulk of the materials, but at the interface between the electrode and the electrolyte. We therefore focused on the interface and succeeded in enhancing the high-rate capability of solid-state lithium batteries by modification of the interface. Thin buffer layers of Li4Ti5O12 [5] or LiNbO3 [6] were interposed between LiCoO2 and the sulfide solid electrolytes, which drastically reduced the interfacial resistance and increased the current drain. The improvement of the high-rate capability was further in the case of LiNbO3, which was attributed to its higher ionic conductivity. In this study, LiTaO3 was used as another buffer layer material, and the electrode performance was compared with those of Li4Ti5O12 and LiNbO3 in order to confirm it, because ionic conductivity of LiTaO3 is higher than Li4Ti5O12 and as high as LiNbO3 [7]. In additions, some details in our previous studies are presented. 2. Experimental
⁎ Corresponding author. National Institute for Materials Science, 1-1, Namiki, Tsukuba, Ibaraki 305-0044, Japan. E-mail address: [email protected] (K. Takada). 0167-2738/$ - see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2008.02.017
In our previous studies, the buffer layers were formed on the surface of LiCoO2 particles in order to interpose them between
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K. Takada et al. / Solid State Ionics 179 (2008) 1333–1337
the LiCoO2 and the sulfide electrolyte, when they were mixed into cathodes. Similarly, in the present study, the LiCoO2 particles were coated with LiTaO3 by spray coating as described below. First, a coating solution of LiTaO3 was prepared. Li metal (1.47 g, 0.21 mol) was dissolved in anhydrous ethanol to obtain an ethanol solution of lithium ethoxide. The solution was mixed with tantalum pentaethoxide (86.26 g, 0.21 mol) and diluted by anhydrous ethanol to make the weight of the final solution 700 g. The solution was sprayed onto the surface of LiCoO2 powder (D10, Toda Kogyo) by a rolling fluidized coating machine (MP-01, Powrex). Different amounts of the solution were sprayed in order to control the thickness of the buffer layer. After the coating, the samples were heated at 400 °C for 30 min under an oxygen flow to decompose the organic components. The samples were labeled by the thickness of the buffer layer, t, in this paper, which was calculated from the amounts of the applied solution and the BET surface area of the LiCoO2 powder (0.26 m2 g− 1). The thickness of the buffer layer was confirmed by transmission electron microscopy. The electrode properties of the samples were investigated in two-electrode cells by the same procedure as described in our previous papers [5,6]. In–Li alloy was used as the counter electrode in the electrochemical measurements, which has an electrode potential of 0.62 V vs. Li+/Li. For clarity, the cell voltages are presented by adding 0.62 V as if lithium metal was used as the counter electrode. Li3.25Ge0.25P0.75S4 was used as the solid electrolyte. Since it is difficult to characterize thin buffer layers on the surface of the LiCoO2 particles, thin films prepared by spin coating and pulse laser deposition (PLD) were characterized instead. X-ray diffraction (XRD) patterns for the spin-coated films were taken in order to investigate the crystal structures of the buffer layers. The ethanol solution of the alkoxides mentioned above and in Refs. [5] and [6] were spin-coated onto Si (100) substrates, and the coated films were then heated at different temperatures for 30 min under oxygen atmosphere. In-plane XRD data were collected for the thin films using synchrotron radiation at Photon Factory BL-3A at the Institute of Materials Structure Science, High Energy Accelerator Research Organization (KEKPF). The wavelength of the incident X-ray was calibrated to be 1.39904(8) Å. An NaI scintillation counter was scanned to take the diffraction data with a grazing exit angle of 0.2° around the normal axis of the thin films. The effect of heat treatment on the buffer layers was investigated using LiCoO2/Li4Ti5O12-stacked films prepared by PLD. LiCoO2 and Li4Ti5O12 powders synthesized by solid-state reactions were sintered into pellets, which were used as targets in the PLD. First, an LiCoO2 film was deposited onto an Si substrate heated at 600 °C using a KrF excimer laser (λ = 248 nm) operated at 10 Hz. Then, an Li4Ti5O12 film was deposited onto the LiCoO2 without heating the substrate to obtain LiCoO2/Li4Ti5O12-stacked films. The stacked films were heated at temperatures ranging from 400 °C to 600 °C for 10 min. Changes in the buffer layer upon the heat treatment were investigated by secondary ion mass spectrometry (SIMS; IMS 4f, CAMACA). Depth profiles for the films heated at different temperatures were measured.
3. Results and discussion 3.1. Strategy for the enhancement of high-rate capability When LiCoO2 and graphite are used as the cathode and the anode, respectively, the rate-determining step is the cathode. As shown in Fig. 1, discharge capacity decreased quickly for LiCoO2 with current density increased, whereas such drastic decrease was not observed for the graphite. In our previous study [5], the reason for the poor high-rate capability of the cathode was considered to be a highly-resistive space-charge layer formed at the interface between the LiCoO2 and the sulfide electrolyte. Electrochemical potentials of the lithium ions should be largely different between oxides and sulfides [8], because oxide ions attract lithium ions much more strongly than sulfide ions. This large difference will cause some of the lithium ions in the sulfide electrolyte to transfer to the LiCoO2 until the equilibrium reached, when they are brought into contact with each other, leaving a lithium-deficient space-charge layer on the electrolyte side at the interface. Since solid electrolytes are optimized in the compositions for fast ionic conduction, the changing lithium content will decrease the ionic conductivity in the space-charge layer. However, contact between oxides and sulfides alone does not always form highly-resistive layer [9]. An even greater problem is electronic conduction in LiCoO2. Interfacial phenomena at the hetero-junction of ionic conductors are now categorized into “nanoionics” [10], in which space-charge layer models for the hetero-junction between ionic conductors are analogous to those at a semiconductor interface. The space-charge layers at the LiCoO2/electrolyte interface will be largely developed as in the Schottky junction because of the electronic conduction in LiCoO2 [11], which will greatly suppress the ionic conduction at the interface. Therefore,
Fig. 1. Discharge curves for LiCoO2 in thio-LISICON (a) [5] and graphite electrodes in a Li2S–P2S5 glass-ceramics electrolyte (b) [4].
K. Takada et al. / Solid State Ionics 179 (2008) 1333–1337
we interposed thin buffer layers of oxides with ionic conduction and electronic insulation, i.e. oxide solid electrolytes, between the LiCoO2 and the solid electrolyte. The interposition forms two interfaces: one between the LiCoO2 and the oxide solid electrolyte, and the other between the oxide and the sulfide electrolyte. The space-charge layers will not be largely developed at both interfaces, because the former interface is between two oxides, between which the electrochemical potential of lithium ions will not be largely different. The latter consists of electronic insulators, where the space-charge layer will be less developed as at p–n junctions between semiconductors [11]. Therefore, we used Li4Ti5O12 as the buffer layer material in our first study [5]. Evidence for the lithium–ion transfer causing large development of the space-charge layer was found in the voltage
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profile at the very beginning of the first charging. Fig. 2 shows change in the voltage profile of the LiCoO2 by the interposition of Li4Ti5O12. When the LiCoO2 was not coated with Li4Ti5O12, the charge curve showed a slope prior to the 4 V plateau. Because some lithium ions should have been transferred from the solid electrolyte to the LiCoO2 during the formation of the space-charge layer, the LiCoO2 should accommodate excess lithium ions. They are considered to be deintercalated in the beginning of the first charging, giving the additional oxidative deintercalation steps appearing as the slopes. The slopes gradually became shorter with increasing amount of applied Li4Ti5O12, which indicates that the increasing coverage of the buffer layer gradually decreased the area where the space-charge layer was developed. Correspondingly, the interfacial resistance appearing as an arc decreased and reached
Fig. 2. Voltage profiles at the beginning of the first charging (left) and complex impedance plots for the cathodes. Dotted, dashed, and solid lines are for Li4Ti5O12, LiNbO3, and LiTaO3 coated LiCoO2, respectively. The complex impedance plots were obtained by subtracting the impedance of the electrolyte layers.
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a minimum at t = 5 nm, although the slope still remained. This means that the buffer layer was so thick that the resistance of the buffer layer itself became predominantly large, when a sufficient amount of Li4Ti5O12 to completely cover the surface was applied. This fact suggests that more conductive buffer layers should be used to improve the high-rate capability. In addition, another requirement for the buffer layers is that they should be formed by low-temperature heat treatment as described below. 3.2. Heat treatment for the formation of the buffer layer Examples of such highly-conductive oxides are Li1 + xMxTi1 − x (PO4)3 with NASICON structure [12] and (Li, La)TiO3 with perovskite structure [13]. However, heat treatment at high temperatures is necessary to attain high ionic conductivity. For instance, LiTi2(PO4)3 thin films prepared by a sol–gel method showed fast ionic conduction only in the crystalline phase after heating above 600 °C [14]. However, buffer layers formed at such high temperatures did not improve the high-rate capability. Fig. 3 compares charge–discharge curves for the LiCoO2 with and without the buffer layer of Li4Ti5O12. When the sample was heated at 600 °C, the buffer layer did not cause any changes in the electrode properties. On the other hand, when the heat treatment temperature was 400 °C, the electrode performance was improved. Changes in the buffer layer upon the heat treatment were investigated by in-plane XRD and SIMS. Fig. 4a shows inplane XRD patterns for the Li4Ti5O12 films prepared by spin coating heated at different temperatures. The film heated at 400 °C did not give any sharp reflections, which suggested the amorphous nature of the film. When the film was heated at 600 °C, sharp reflections appeared, indicating crystallization of the film. That is, the buffer layer became ineffective by the heat treatment at 600 °C, although it was crystallized.
Fig. 3. Charge–discharge curves for uncoated (upper) and Li4Ti5O12-coated LiCoO2 heated at 400 °C and 600 °C at a current density of 127 μA cm− 2.
Fig. 4. Synchrotron radiation in-plane diffraction patterns for the thin films prepared on Si (100) substrates by spin-coating. (a) Patterns for Li4Ti5O12 heated at different temperatures. The reflections were indexed on a cubic lattice with a lattice constant of a = 8.311(5) Å. (b) Pattern for LiTaO3 prepared at 400 °C. The reflections were indexed on a rhombohedral lattice with lattice constants of a = 5.142(2) Å and c = 13.87(1) Å.
The reason can be seen in the depth profiles obtained for the LiCoO2/Li4Ti5O12-stacked films shown in Fig. 5. These depth profiles clearly indicated that the heat treatment at high temperature caused cobalt ions to diffuse into the Li4Ti5O12 layer, even when the heating duration was only 10 min and the thickness of the Li4Ti5O12 layer was ca. 150 nm. The cobalt ions penetrating the Li4Ti5O12 layer should induce electronic conduction in it, and thus the buffer layer can not adequately suppress the development of the space-charge layer. Therefore, we used LiNbO3 in place of Li4Ti5O12 to improve the high-rate capability [6]. Because LiNbO3 has high ionic conduction in the
Fig. 5. Depth profiles of LiCoO2/Li4Ti5O12-stacked films heated at different temperatures. The vertical axis indicates the intensity ratio of Co/6Li measured by SIMS.
K. Takada et al. / Solid State Ionics 179 (2008) 1333–1337
amorphous state [7], the heating temperature for the buffer layer formation can be reduced. LiTaO3 was also reported to have fast ionic conduction in the amorphous state as LiNbO3, and thus LiTaO3 was used as the buffer layer in the present study. A thin film of LiTaO3 prepared by spin coating and heated at 400 °C gave reflections in the inplane XRD pattern as shown in Fig. 4a. However, halos were also observed, indicating that at least part of the film remained amorphous. 3.3. Improvement of high-rate capability by the interposition of lithium tantalum oxide buffer layer Electrode properties of the LiTaO3-coated LiCoO2 are presented with those of the Li4Ti5O12 and LiNbO3-coated one in Fig. 2. The slopes, which correspond to the capacities of the space-charge layers, were in good agreement with those observed for Li4Ti5O12 and LiNbO3. That is, the change depended only on the applied amount of buffer layer materials, and not on the kind of materials. If the increasing thickness is the reason for the gradual change, different buffer materials should cause different changes. Therefore, it is concluded again that the increasing coverage is responsible for the gradual shortening of the slope. On the other hand, the LiTaO3 reduced the interfacial resistance more prominently than the Li4Ti5O12 and equivalently to the LiNbO3; the minimum interfacial resistance was almost the same as that observed for the LiNbO3 system. 4. Conclusions As presented in this paper, oxide solid electrolytes introduced to solid-state batteries with sulfide electrolytes as buffer layers improved the high-rate capability; that is, two kinds of solid electrolytes were used to enhance the performance. In our previous study on improving energy density of solid-state lithium batteries [3], we used two kinds of solid electrolytes in a battery: one resistive to highly-oxidizing
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LiCoO2 and the other resistive to highly-reducing C6Lix. These studies may be categorized as a “multi-electrolyte system”, which is valid only in solid-state batteries and effective for improving battery performance. Acknowledgements The authors thank T. Kurihara of Powrex Corp. for his help in preparing the coated LiCoO2. This work was partially supported by the Ministry of Economy, Trade, and Industry (MITI), New Energy, Industrial Technology Development Organization (NEDO), and the Grant-in-Aid for Scientific Research on Priority Area, “Nanoionics (439)” (19017018) by the Ministry of Education, Culture, Sports, Science and Technology (MEXT) of Japan. References [1] R. Kanno, M. Murayama, J. Electrochem. Soc. 148 (2001) A742. [2] F. Mizuno, A. Hayashi, K. Tadanaga, M. Tatsumisago, Adv. Mater. 17 (2005) 918. [3] K. Takada, T. Inada, A. Kajiyama, H. Sasaki, S. Kondo, M. Watanabe, M. Murayama, R. Kanno, Solid State Ionics 158 (2003) 269. [4] Y. Seino, K. Takada, B.-C. Kim, L. Zhang, N. Ohta, H. Wada, M. Osada, T. Sasaki, Solid State Ionics 176 (2005) 2389. [5] N. Ohta, K. Takada, L. Zhang, R. Ma, M. Osada, T. Sasaki, Adv. Mater. 18 (2006) 2226. [6] N. Ohta, K. Takada, I. Sakaguchi, L. Zhang, R. Ma, K. Fukuda, M. Osada, T. Sasaki, Electrochem. Commun. 9 (2007) 1486. [7] A.M. Glass, K. Nassau, T.J. Negran, J. Appl. Phys. 49 (1978) 4808. [8] P.G. Bruce, Chem. Commun. (1997) 1817. [9] K. Takada, M. Tansho, I. Yanase, T. Inada, A. Kajiyama, M. Kouguchi, S. Kondo, M. Watanabe, Solid State Ionics 139 (2001) 241. [10] J. Maier, Nat. Matters 4 (2005) 805. [11] J. Maier, Prog. Solid State Chem. 23 (1995) 171. [12] H. Aono, E. Sugimoto, Y. Sadaoka, N. Imanaka, G. Adachi, J. Electrochem. Soc. 136 (1989) 590. [13] Y. Inaguma, L. Chen, M. Itoh, T. Nakamura, T. Uchida, M. Ikuta, M. Wakihara, Solid State Commun. 86 (1993) 689. [14] K. Takada, K. Fujimoto, T. Inada, A. Kajiyama, M. Kouguchi, S. Kondo, M. Watanabe, Appl. Surf. Sci. 189 (2002) 300.