Schematic representation of Li growth during DC polarization experiments in the investigated garnet-type metal oxides.

Li-stuffed garnet-type solid Li-ion electrolytes have been considered as a promising candidate for all-solid-state Li batteries. In this work, alkaline earth metal-doped Li garnet-type solid Li-ion electrolytes, Li6.5La2.5A0.5TaZrO12 (A = Ca, Sr, Ba) were prepared by conventional solid-state sintering (CSS) and spark plasma sintering (SPS) methods. The effect of sintering methods on the structural and electrochemical properties of the solid electrolytes was investigated by powder X-ray diffraction (PXRD), scanning electron microscopy (SEM), and electrochemical impedance spectroscopy (EIS). Among the investigated compounds, Sr-doped garnet-type Li6.5La2.5Sr0.5TaZrO12 prepared by the SPS method showed the highest Li-ion conductivity of 3.08 × 10–4 S cm–1 at 20 °C and the lowest activation energy of 0.35 eV. Both Sr- and Ba-doped samples exhibited a critical current density of 0.15 mA cm–2, and the Sr-doped Li6.5La2.5Sr0.5TaZrO12 sample showed the lowest Li-ion charge-transfer resistance of 139 Ω cm2 at room temperature.

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Garnet-like lithium-ion conducting solid electrolytes attracts great interests for their potential application to all solid state batteries with lithium anode. One of drawbacks of this material is severer sintering condition (typically >1200°C and > 10 h). It has been demonstrated that spark plasma sintering (SPS) successfully accelerates the sintering of garnet-like solid electrolytes like other materials. On SPS, large pulse current flows across grain interfaces to cause high temperature locally and facilitate sintering. It is possible that such locally generated heat would cause inhomogeneity in the obtained specimen. In this paper, we prepared pellets of garnet-like Li6.5La3Zr1.5Ta0.5O12 (LLZT) and investigated the sintering mechanism by focusing on the inhomogeneity inside the pellets. In addition, the ionic conductivity as well as anti-short-circuit properties will be also demonstrated. LLZT powder was synthesized by solid state reactions using LiOH, La2O3, ZrO2 and Ta2O5. Except for Li, other starting materials were mixed with a stoichiometric ratio. Amount of Li source was by 10% higher than the stoichiometric ratio to compensate the loss of Li on the following calcination, which was conducted at 900°C for 10 h in air. Calcined powder was sintered using graphite dies by SPS with a sintering temperature of 900-1100°C, pressure of 12.5-50 MPa and time of 1-30 min under vacuum. Crystalline structure of LLZT powder and sintered pellets were confirmed by XRD. To confirm distribution of phases along depth, XRD profiles were repeatedly recorded after surface of the pellets was polished for certain thickness. On one side of pellets, which corresponds to anode of pulse current of SPS, La2Zr2O7 was observed. This impurity was confirmed only on the anode and disappeared with polishing the surface, suggesting inhomogeneous decomposition, possibly caused by electrolysis. Interestingly, on cathode, no impurity phase was confirmed by XRD. When the La2Zr2O7 is formed as a result of electrolysis, reduction should occur on cathode. The cathode product is possibly amorphous carbon reduced by Li2CO3 that had been produced on surface of LLZT particles by absorption of CO2. The Li2CO3 is supposed to play another role on SPS. When LLZT powder with smaller particle size was used for SPS, more La2Zr2O7 was produced, suggesting the resistance between electrodes were reduced by molten Li2CO3. In addition to the La2Zr2O7, asymmetric impurity formation was visually confirmed. On anode side, the pellets were white, while cathode sides looked dark grey to black. However, the XRD did not show any impurity phases, indicating the origin of the dark color was amorphous. By taking account of the electrolysis of LLZT on anodes to form La2Zr2O7, it is plausible that amorphous carbon was produced on cathodes due to the electrochemical reduction of CO3 2− of Li2CO3. These impurities were formed in the region within ca. 100-300 μm from the surface of pellets for both cathode and anode sides. Therefore, it is necessary to polish pellets for at least 300 μm for both sides before the pellets are used for electrochemical cells. When these impurities were removed, the obtained pellets exhibited very high density of 96% and high ionic conductivity of 7×10−4 S cm−1 at 298 K, which are attractive properties for all solid state batteries. DC polarization was applied to investigate the anti-short-circuit properties of the pellets using a symmetric cell of Li | LLZT | Li. Direct current density was progressively increased from ±20 μA cm−2. Each dc polarization was applied for 30 min. for one direction, and after repeating 5 cycles for both directions, next higher current density was applied. Although short-circuit was confirmed at 200 μA cm−2 (Figure), it was confirmed that the resistance gradually decreased, indicating the dendrite of Li was growing inside the LLZT even at low current density of 20 μA cm−2. The limit current density of 200 μA cm−2 was comparable to reported values on pellets prepared by conventional solid state sintering process. Figure 1

Recently, lithium ion conducting solid electrolytes attract many interests because of their potential application to all-solid-state batteries (ASSBs), which are supposed to be one of candidates of next-generation energy storage devices with high safety and reliability [1-3]. With great efforts of development of solid electrolytes, several materials with high lithium ionic conductivity have been reported, which facilitated hope of realization of ASSBs [2,4-8]. Now, researchers face several problems on application of solid electrolytes to ASSBs. Most of those problems are related to interfaces of between solid electrolytes and active materials, for example, contact areas, interdiffusion of elements on assembly and/or electrochemical processes of ASSBs. To solve these issues, much attention has been paid to design and fabrication techniques. In addition to the problems on the fabrication, several reports point out intrinsic phenomena related change in ion concentration at interfaces of solid electrolytes [1, 9, 10]. We have so far studied the local structure at the interfacial region by fabricating nano-composites of active materials and solid electrolytes, and confirmed that the interfacial region with distorted lattice ranges for at least 20-50 nm in thickness [9, 10]. Although more researches with variety of experimental analysis are required for further comprehension of the interfacial phenomena, difficulty in the analysis of solid/solid interface limits experimental tools. It is known that surface of solid electrolytes is also related to interfacial resistance. Kliewer proposed cation depletion at surface of ionic crystals when free energy of defect formation is lower for cation than for anion [11]. In addition, we observed that ball-milling of a poor ionic conductor (Li2SiO3, LSO) resulted in increase in ionic conductivity without change in activation energy [9, 12]. In this study, three different surface-sensitive techniques, grazing incidence X-ray diffraction (GIXD), attenuated total reflectance Fourier transform infrared (ATR-FTIR), and X-ray photoelectron spectroscopy (XPS), were employed to analyze the local structure and composition around the surface region of a lithium ion conducting solid electrolyte sheet (NASICON-type in the system of Li-Al-Ti-Ge-Si-P-O, Ohara Inc. abbreviated as LICGC). It was revealed that the local structure and composition changed depending on depth from the top surface of the sheet. At the very surface, there was a layer with expanded lattice (Figure 1) and high Li composition (Figure 2). And with increasing the depth from the surface, lattice shrank sharply, and then expanded again gradually. The change in the lattice seemed to be accompanied by Li and Al composition. It was supposed that the change in the Li composition and structure is induced by combination of chemical reaction, segregation, and defects distribution [13]. References [1] N. Ohta, K. Takada, L. Zhang, R. Ma, M. Osada, T. Sasaki, Adv. Mater. 18 (2006) 2226-2229. [2] T. Kobayashi, Y. Imade, D. Shishihara, K. Homma, M. Nagao, R. Watanabe, T. Yokoi, A. Yamada, R. Kanno, T. Tatsumi, J. Power Sources, 182 (2008) 621-625. [3] M. Kotobuki, H. Munakata, K. Kanamura, Y. Sato, T. Yoshida, J. Electrochem. Soc. 157 (2010) A1076-A1079. [4] P. Knauth, Solid State Ionics 180 (2009) 911-916. [5] G. Adachi, N. Imanaka, H. Aono, Adv. Mater. 8 (1996) 127-135. [6] Y. Inaguma, L. Chen, M. Itoh, T. Nakamura, Solid State Commun. 86 (1993) 689-693. [7] M. Tatsumisago, A. Hayashi, Solid State Ionics 225 (2012) 342-345. [8] N. Kamaya, K. Homma, Y. Yamakawa, M. Hirayama, R. Kanno, M. Yonemura, T. Kamiyama, Y. Kato, S. Hama, K. Kawamoto, A. Mitsui, Nature Mater. 10 (2011) 682-686. [9] H. Yamada, Y. Oga, I. Saruwatari, I. Moriguchi, J. Electrochem. Soc. 159 (2012) A380-A385. [10] H. Yamada, K. Suzuki, K. Nishio, K. Takemoto, G. Isomichi, I. Moriguchi, Solid State Ionics 262 (2014) 879-882. [11] K. L. Kliewer, J. S. Köhler, Phys. Rev. 140 (1965) A1226-A1240. [12] H. Yamada, D. Tsunoe, S. Shiraishi, J. Phys. Chem. C 119 (2015) 5412-5419. [13] H. Yamada, K. Takemoto, Solid State Ionics 285 (2016) 41-46. Figure 1