Development of Sulfide Solid Electrolytes and Interface Formation Processes for Bulk-Type All-Solid-State Li and Na Batteries

Introduction

All-solid-state batteries using inorganic solid electrolytes (SEs), used in place of conventional organic liquid electrolytes, have been studied because of their high safety (non-flammability with no liquid leakage), long cycle life, and versatile geometries (Takada, 2013; Tatsumisago et al., 2013; Tatsumisago and Hayashi, 2014). These features are important for large rechargeable lithium batteries with high energy density for application to eco-cars, such as electric vehicles and plug-in hybrid vehicles. Rechargeable sodium batteries are also attractive for large-scale applications for stationary load-leveling because sodium is expected to be the next targeted element after lithium based on its atomic weight, standard potential, and natural abundance (Yabuchi et al., 2015b; Yamada, 2014).

A key material to realize all-solid-state rechargeable batteries is a superior SE. In lithium ion conductors, sulfide electrolytes of Li₁₀GeP₂S₁₂ and Li₇P₃S₁₁ have high room-temperature conductivity of more than 10⁻² S cm⁻¹, which is as high as the conductivity of conventional organic liquid electrolytes (Kamaya et al., 2011; Seino et al., 2014). The lithium ion transport number of liquid electrolytes is below 0.5, whereas that of SE is 1. It is noteworthy that the conductivity of lithium ions of sulfide SEs already exceeds that of organic liquid electrolytes. Very recently, Li_9.54Si_1.74P_1.44S_11.7Cl_0.3 has been reported to show the highest conductivity of 2.5 × 10⁻² S cm⁻¹. Using this superior electrolyte, high power competing with that of supercapacitors can be achieved in all-solid-state rechargeable lithium batteries (Kato et al., 2016b). Sulfide SEs have several important benefits of high conductivity, wide electrochemical window, and appropriate mechanical properties, such as formability and elastic modulus. A shortcoming of sulfide electrolytes is its low air stability. To realize bulk-type all-solid-state batteries, the formation of favorable solid–solid interfaces between electrode and electrolyte is important in addition to the development of superior sulfide electrolytes. Direct coating of sulfide electrolytes on active material particles instead of adding electrolyte particle is effective for forming a close solid–solid interface with large contact area. Insulative active materials, such as sulfur and Li₂S, should be blended not only with SE but also with carbon-conductive additive to form ion and electron conduction paths to active materials. Nanocomposites of the three components are useful by preparation with high-energy ball milling, which pulverizes and combines them. Transition metal sulfides (MS_x) with high conductivity have good compatibility with sulfide SEs having the same sulfide anions. Especially, sulfur-rich compounds, such as TiS₃, are attractive active materials with high reversible capacity. Lithium metal is a supremely negative electrode, but issues related to lithium dendrites prevent its commercialization. Combination with a SE is a promising solution. Interface modification between Li metal and the SE is important.

As described in this paper, the recent development of sulfide SEs and interface formation processes for bulk-type all-solid-state Li and Na batteries are reviewed. Significant properties as SEs of conductivity, chemical stability, and mechanical property for Li⁺ or Na⁺ ion conducting sulfides are reported. Procedures for preparing sulfide electrolytes, such as mechanical milling and liquid-phase synthesis, are also described. Several approaches to form favorable solid–solid interfaces developed by our research group are demonstrated. Processes for the coating of sulfide electrolytes via gas-phase or liquid-phase process on LiCoO₂ or graphite particles have been developed. Preparation of nanocomposite electrodes with sulfur, Li₂S, and Li₃PS₄ (as a bifunctional material of electrolyte and electrode) particles is described, along with the use of MS_x active materials. Formation of a solid–solid interface for using lithium metal negative electrodes is also discussed.

Development of Sulfide Solid Electrolytes

Sulfide SEs with Li⁺ or Na⁺ ion conductivity have been developed during the past three decades. Recently, chemical stability and mechanical properties as well as conductivity for sulfide electrolytes have attracted much attention. Sulfide electrolytes are prepared using several techniques with solid-phase reaction, melt-quenching, mechanical milling, crystallization of mother glasses, and liquid-phase reaction. Detailed information related to electrolyte properties will be presented in the following sections of this report.

Conductivity

Inorganic sulfide SEs with high Li⁺ or Na⁺ ion conductivities have been developed. Sulfide electrolytes with high conductivities are presented in Table 1. In Li⁺ ion conducting sulfide electrolytes, crystalline Li₁₀GeP₂S₁₂ (LGPS; Kamaya et al., 2011) and glass–ceramic Li₇P₃S₁₁ (Seino et al., 2014) have considerably high conductivity of more than 10⁻² S cm⁻¹ at 25°C, which is beyond the Li⁺ ion conductivity of conventional organic liquid electrolytes. Very recently, Li_9.54Si_1.74P_1.44S_11.7Cl_0.3 with LGPS structure has been reported to show the highest conductivity of 2.5 × 10⁻² S cm⁻¹ (Kato et al., 2016b). Studies investigating new electrolytes with much higher conductivity are in progress. Several crystals, such as Li₁₀SnP₂S₁₂ (Boron et al., 2013) and Li₆PS₅Cl (Boulineau et al., 2012), have high conductivity of more than 10⁻³ S cm⁻¹. This conductivity is also achieved by the addition of lithium halides to sulfide glass and glass–ceramic electrolytes (Wada et al., 1983; Ujiie et al., 2014).

TABLE 1

Table 1. Li⁺ ion and Na⁺ ion conductivities for sulfide solid electrolytes.

The Na⁺ ion conductivity is lower than Li⁺ ion conductivity in glassy electrolytes. Ionic conduction of Na⁺ ion with ionic radius larger than Li⁺ ion is unfavorable in glasses (Souquet et al., 1981). For sodium ion conductors, sulfides with high conductivity had not been found since this report described Na₃PS₄ glass–ceramic electrolytes in 2012. A cubic Na₃PS₄ phase is precipitated by crystallization of a mother Na₃PS₄ glass. The prepared glass–ceramic electrolyte shows Na⁺ ion conductivity of greater than 10⁻⁴ S cm⁻¹ at 25°C (Hayashi et al., 2012b). Furthermore, partial substitution of Si for P in Na₃PS₄ is useful for increasing conductivity (Tanibata et al., 2014). Figure 1 presents the composition dependence of conductivities at 25°C for Na₃PS₄–Na₄SiS₄ glass–ceramic electrolytes. The replacement of 6 mol% Na₃PS₄ by Na₄SiS₄ increases the conductivity from 4.6 × 10⁻⁴ to 7.4 × 10⁻⁴ S cm⁻¹. The electron density distribution of the cubic Na₃PS₄ structure obtained using the maximum entropy method is shown in the inset of Figure 1. Na₃PS₄ has three-dimensional Na⁺ ion conduction paths along Na1 and Na2 sites. Increased Na⁺ ion concentration is presumed as the reason for the enhancement of conductivity. Very recently, high conductivity of more than 10⁻³ S cm⁻¹ has been reported in crystalline Na₃PSe₄ (Zhang et al., 2015). However, studies of new sulfide Na⁺ ion conductors are few. A first-principles calculation indicates that Sn-doped cubic Na₃PS₄ is predicted to have a higher Na⁺ ion conductivity of 10⁻² S cm⁻¹ (Zhu et al., 2015). Synthesis of SEs with a favorable structure for Na⁺ ion conduction based on calculation results is important for finding new Na⁺ ion conducting SEs.

FIGURE 1

Figure 1. Composition dependence of conductivity at 25°C for (100-x)Na₃PS₄·xNa₄SiS₄ (mol%) glass–ceramic electrolytes. The inset shows the electron density distribution of the cubic Na₃PS₄ structure.

Chemical Stability

A shortcoming of sulfide electrolytes is their lower chemical stability in air atmosphere. Sulfides tend to be decomposed by hydrolysis, generating harmful H₂S. Suppression of hydrolysis of sulfides is an important task for developing sulfide electrolytes. Based on our early experiments, the selection of compositions in sulfide electrolytes gives moderate stability in air to sulfide electrolytes (Muramatsu et al., 2011). The chemical stability of sulfide glass electrolytes in the binary system Li₂S–P₂S₅ was examined by exposing them to air atmosphere. Amounts of generated H₂S from the sulfides depend on the composition of the glasses; H₂S generation is minimized at the composition 75Li₂S⋅25P₂S₅ (mol%). The glass comprises Li⁺ and ${\text{PS}}_4^{3-}$ ions. An isolated anion ${\text{PS}}_4^{3-}$ without bridging sulfurs is useful for high tolerance for hydrolysis. Li₃PS₄-based SEs with both good chemical stability and high conductivity have been prepared by the combination of oxides (Li₂O or P₂O₅) and iodides (LiI) (Ohtomo et al., 2013a,b). The addition of metal oxides, such as ZnO, which act as an absorbent for H₂S, is also effective for decreasing H₂S (Hayashi et al., 2013). It is noteworthy that the use of a favorable M_xO_y (M_xO_y: Fe₂O₃, ZnO, and Bi₂O₃) with a larger negative Gibbs energy change (ΔG) for the reaction with H₂S is effective for improving the chemical stability of sulfide electrolytes. Another approach is the use of sulfide compositions based on the hard and soft acids and bases theory (HSAB; Sahu et al., 2014). Lithium tin thiophosphate, Li₄SnS₄, has better air stability than that of Li₃PS₄. Actually, as-substituted Li₄SnS₄ has good features of both high conductivity of 10⁻³ S cm⁻¹ and high air stability.

The chemical stability of SEs tends to affect battery performance. The electrochemical performance of all-solid-state C/LiCoO₂ cells using Li₃PS₄ glass or Li₇P₃S₁₁ glass–ceramic as a SE is compared. The cell with Li₃PS₄ glass electrolytes exhibits better cycle performance, although Li₃PS₄ glass has lower conductivity than Li₇P₃S₁₁ electrolyte (Ohtomo et al., 2013c). High performance of the battery would be based on the higher chemical stability of Li₃PS₄ electrolytes. Chemical stability as well as conductivity is an important factor of SEs for developing superior solid-state batteries.

Mechanical Property

Adhesion of the solid–solid interface is a key to the utilization of electrode active materials in all-solid-state batteries. Formability or processability of SEs is examined by the molding pressure dependence of the relative density of compressed powder pellets.

Figure 2A shows the dependence of the relative density of 75Li₂S⋅25P₂S₅ and 75Na₂S⋅25P₂S₅ (mol%) glass electrolytes on molding pressure. The relative densities increase gradually with an increase in molding pressure in both glasses at the same alkali compositions of 75 mol% M₂S (M = Li or Na), whereas the relative density for the Na₂S system is higher than that for the Li₂S system (Sakuda et al., 2013a,b; Nose et al., 2015). Cross-sectional SEM images of the 75 mol% M₂S pellets pressed at 360 MPa are shown in the inset. Grain boundaries and voids in the pellets are more decreased in the 75Na₂S⋅25P₂S₅ glass compared with the 75Li₂S⋅25P₂S₅ glass. Sulfide glasses are densified by cold pressing without heat treatment, and this densification phenomenon is called “room-temperature pressure sintering” (Sakuda et al., 2013a,b). It is noteworthy that the 75Na₂S⋅25P₂S₅ glass has better formability than the 75Li₂S⋅25P₂S₅ glass. Both the glasses comprise Li⁺ or Na⁺ ion and ${\text{PS}}_4^{3-}$ ion, which are thought to diffuse at the particle boundaries on pressing at room temperature. Na⁺ ion with a larger ionic radius than Li⁺ ion has a weaker interaction with ${\text{PS}}_4^{3-}$ ion. Therefore, both Na⁺ and ${\text{PS}}_4^{3-}$ ions would diffuse readily by cold press, leading to better densification.

FIGURE 2

Figure 2. (A) Molding pressure dependence of relative density of 75M₂S⋅25P₂S₅ (M = Li or Na, mol%) glass electrolyte. Cross-sectional SEM images of these glass pellets prepared by pressing at 360 MPa are also shown in the inset. (B) Composition dependence of Young’s modulus for the xM₂S⋅(100-x)P₂S₅ glasses.

Retaining solid–solid contacts between active materials and SEs during charge–discharge processes brings about long cycle lives of all-solid-state batteries. Young’s moduli of SEs are important for keeping favorable contacts even at volume changes of active materials. Those for densified sulfide electrolytes prepared by hot-pressing are determined by an ultrasonic pulse-echo technique and the uniaxial compression tests (Sakuda et al., 2013a,b; Nose et al., 2015). Young’s moduli of the sulfide glasses in the systems Li₂S–P₂S₅ and Na₂S–P₂S₅ are presented in Figure 2B. They are increased gradually with the increase in the alkali content in both systems. The Na₂S–P₂S₅ glasses have smaller Young’s moduli of 15–19 GPa than the Li₂S–P₂S₅ glasses (18–25 GPa). The difference on Young’s modulus is understood based on the Coulomb force and the mean atomic volume of the glasses. These sulfide glasses have an intermediate Young’s modulus between oxide glasses and organic polymers. Sulfide electrolytes deforming elastically are expected to act as a buffer in response to volume changes of active materials during charge–discharge processes. In fact, most all-solid-state batteries that use sulfide SEs exhibit good cycle performance.

Preparation Process

Sulfide SEs were prepared via various techniques. Crystalline electrolytes are prepared via solid-phase reaction, whereas glass electrolytes are obtained using the melt-quenching method. In general, sulfide starting materials are sealed in a carbon-coated quartz tube under vacuum and then heat-treated because of high vapor pressure of sulfides at high temperatures. Heating temperatures and cooling rates at preparation process affect precipitated crystals. A phase diagram in ternary system Li₂S–GeS₂–P₂S₅ is complicated (Hori et al., 2015), and crystalline Li₁₀GeP₂S₁₂ with a conductivity of 10⁻² S cm⁻¹ is prepared by selecting experimental conditions. Another preparation technique of sulfide electrolytes is mechanical milling using a high-energy planetary ball mill apparatus. This technique is fundamentally a room-temperature process. Therefore, sulfides are reacted at ordinary temperature and pressure. Electrolyte particles are obtained directly by milling. They are applicable to all-solid-state batteries without additional pulverization of electrolytes. Crystallization of glass electrolyte tends to precipitate a metastable phase, such as high-temperature phase, which generally has high ionic conductivity (Tatsumisago et al., 1991; Hayashi et al., 2003). Crystalline Li₇P₃S₁₁ (a high-temperature phase at the composition) is precipitated as a primary phase by heat treatment of a corresponding mother glass, and the obtained glass–ceramic electrolytes show conductivities of 10⁻³–10⁻² S cm⁻¹, which depend on the degree of crystallinity and the grain boundary of Li₇P₃S₁₁ (Mizuno et al., 2005, 2006; Seino et al., 2014). Cubic Na₃PS₄ is crystallized from the Na₃PS₄ glass prepared by mechanical milling. The prepared glass–ceramic electrolytes with Na₃PS₄ show Na⁺ ion conductivity of 10⁻⁴ S cm⁻¹ at 25°C (Hayashi et al., 2012b).

Electrolyte preparation via liquid phase is a suitable process for cost-effective quantity synthesis without using a special reaction apparatus. In general, this process has benefits of lowering the reaction temperature, shortening the reaction time, and controlling the particle morphology and size. Prepared electrolyte solutions are also useful for the coating of active material particles. Very recently, liquid-phase synthesis of sulfide SEs has been reported. The reaction processes for sulfides are divided into two categories: one uses suspension and the other uses a homogeneous solution. For the former synthesis, β-Li₃PS₄ synthesized in tetrahydrofuran (Liu et al., 2013) or dimethyl carbonate (Phuc et al., 2016) as reaction medium and Li₇P₃S₁₁ synthesized in 1,2-dimethoxyethane (Ito et al., 2014) are reported. In these reactions, precursors with precipitates are obtained. Compressed pellets of the heat-treated sulfide electrolytes show conductivity of greater than 10⁻⁴ S cm⁻¹ at 25°C. As the latter one via homogeneous liquid, Li₃PS₄ is synthesized from a mixture of Li₂S and P₂S₅ with N-methlyformamide (NMF) (Teragawa et al., 2014a). The Li₃PS₄ SEs can also be prepared using a dissolution–reprecipitation process in NMF from 80Li₂S⋅20P₂S₅ (mol%) glass prepared in advance using mechanical milling (Teragawa et al., 2014b). The prepared Li₃PS₄ electrolyte shows low conductivity of 10⁻⁶ S cm⁻¹ at the present stage, but conductivity can be enhanced by selecting electrolyte compositions. Argyrodide-type Li₆PS₅Cl is dissolved into ethanol. Then, the argyrodite phase is reprecipitated by removing ethanol at 80°C under vacuum for 3 h (Yubuchi et al., 2015a,b). A pellet of the product shows conductivity of 10⁻⁵ S cm⁻¹ at 25°C, which is somewhat lower than that of the original Li₆PS₅Cl. Grain boundary resistance, which is affected by surface structure and morphology, might be greater in the prepared Li₆PS₅Cl. Optimization of posttreatments for the prepared powders will enhance the Li₆PS₅Cl conductivity. Furthermore, the combination of sulfide SEs and ionic liquids produces pseudo-SEs (Minami et al., 2010; Oh et al., 2015). The prepared electrolytes give a new category of electrolytes having both high conductivity and good formability.

Sodium-ion conducting sulfide electrolytes with cubic Na₃PS₄ are also synthesized via a liquid-phase process from the mixture of Na₂S and P₂S₅ in NMF solvent (Yubuchi et al., 2015a). The room-temperature conductivity of the obtained electrolyte is 10⁻⁶ S cm⁻¹, which is lower than the conductivity of the electrolyte prepared by mechanical milling. Studies investigating sulfide SEs with Na⁺ ion conductivity are extremely few at present. New electrolytes produced via a simple liquid-phase process will be researched widely.

Preparation of Solid–Solid Interface in All-Solid-State Batteries

Sulfide glasses are well-balanced SEs with high conductivity, good formability, appropriate Young’s modulus, and moderate chemical stability. They are therefore highly promising SEs for use in all-solid-state batteries. A schematic diagram of bulk-type all-solid-state batteries is depicted in Figure 3. A positive electrode layer is composed not only of active material particles but also of SE ones. The Li⁺ ion is supplied from SEs attached to active materials. Electrons are mobile through active materials. To enhance the rate of performance of the batteries, conductive additives of nanocarbons are added to the electrode layer. A lithium alloy or lithium metal is used as the negative electrode. A SE layer as a separator is sandwiched with the positive and negative electrode layers. Then, because of good formability of sulfide SEs, it can be pressed uniaxially at room temperature to fabricate bulk-type all-solid-state batteries.

FIGURE 3

Figure 3. Schematic diagram of bulk-type all-solid-state batteries.

An all-solid-state In/LiCoO₂ or Li-In/Li₄Ti₅O₁₂ cell with Li₂S–P₂S₅ glass–ceramic electrolytes exhibits good cycle performance for hundreds of times at 25°C (Tatsumisago and Hayashi, 2008; Tatsumisago et al., 2013). These cells operate at a high temperature of 100°C, where it is difficult for a liquid electrolyte cell to be used. The all-solid-state cell with Li₄Ti₅O₁₂ shows a discharge–charge capacity of about 140 mAh g⁻¹. It retains the capacity for 700 cycles with no degradation under a high current density of more than 10 mA cm⁻² (Minami et al., 2011).

Coating of Sulfide Electrolytes on LiCoO₂ and Graphite Particles

Direct coating of sulfide SEs on LiCoO₂ particles, instead of mixing electrolyte particles, is effective for forming good electrolyte–LiCoO₂ interfaces with a wide contact area. Pulsed laser deposition (PLD) was first used as a coating technique (Sakuda et al., 2010, 2011). In this study, LiNbO₃-coated LiCoO₂ particles were fluidized with a vibrator to ensure uniform coating of sulfide electrolytes on the particles.

Figure 4 presents cross-sectional SEM images of positive electrodes with (a) a conventional mixture of LiCoO₂ and Li₂S–P₂S₅ SE particles with the weight ratio of 70/30 and (b) SE-coated LiCoO₂ particles, where the weight ratio of LiCoO₂ and SE coatings was 90/10. Figure 2A shows that solid–solid contacts between LiCoO₂ and SE are formed by cold pressing of the mixture electrode because SE particles are deformed easily by pressing. Figure 4A shows that the aggregation of LiCoO₂ particles engenders less utilization of LiCoO₂ because many voids and less contact with SE are observed among LiCoO₂ particles. A close-packed electrode layer is formed using only SE-coated LiCoO₂ (Figure 4B). Its enlarged figure, which will be displayed in Figure 7D later, shows that a close solid–solid interface with fewer voids appears, and Li⁺ ion conduction paths are therefore formed among LiCoO₂ particles. It is noteworthy that SE amounts in the electrode are decreased considerably by SE coating on LiCoO₂ particles. The use of minimum amounts of SE in an electrode layer contributes to enhanced energy density of all-solid-state batteries.

FIGURE 4

Figure 4. Cross-sectional SEM images of LiCoO₂ positive electrode layers consisting of (A) mixture electrode with LiCoO₂ and Li₂S–P₂S₅ electrolyte particles and (B) electrolyte-coated LiCoO₂ particles.

A cross-sectional high-angle annular dark field (HAADF)-STEM image and EDX mappings of O, P, S, Co, and Nb elements for the SE-coated LiCoO₂ electrodes are presented in Figure 5. A SE coating layer is observed between two LiCoO₂ particles, where one particle has several cracks. The EDX mapping of Nb element reveals that a LiNbO₃ coating layer exists on the surface of LiCoO₂, but the layer is missing at the cracks. The EDX mappings of P and S elements indicate that SE penetrates into the cracks as a liquid electrolyte does. Favorable formability of sulfide SE is effective for forming good contacts with the surface of active materials and using active materials even at newly formed crystal faces with cracks.

FIGURE 5

Figure 5. Cross-sectional high-angle annular dark field (HAADF)-TEM image and EDX mappings of O, P, S, Co, and Nb elements for the SE-coated LiCoO₂ electrodes.

Solid electrolyte coatings were also done for graphite particles using PLD. The mixture of SE-coated graphite and SE particles was used as a negative electrode. The weight ratio of graphite and SE was 90/10. All-solid-state cells with Li₂S–P₂S₅ SE as a separator layer were charged and discharged at a constant current density of 0.064 mA cm⁻² (ca. 0.05 C) for voltages of 2.8–4.3 or 4.6 V at room temperature (Sakuda et al., 2013b). Figure 6 shows charge–discharge curves for all-solid-state SE-coated graphite/SE-coated LiCoO₂ batteries. The battery with a cutoff voltage of 4.6 V at the charge process shows a higher discharge capacity than that of 4.3 V. Battery has a discharge capacity of 133 mAh g^–1, as calculated from the total mass of the composite positive electrode.

FIGURE 6

Figure 6. Initial charge–discharge curves for an all-solid-state SE-coated graphite/SE-coated LiCoO₂ battery.

Figure 7 shows cross-sectional SEM images of SE-coated LiCoO₂ positive electrodes; Figure 7A as-prepared, Figure 7B after the initial charge-discharge (cutoff voltage: 4.3 V), and Figure 7C after the initial charge–discharge (cutoff voltage: 4.6 V). Figures 7D–F, respectively, portray enlarged images of Figures 7A–C. Closely attached SE–LiCoO₂ interfaces with large contact area are achieved by PLD coating, as shown in Figures 7A,D. The electrode morphology (Figure 7B) does not change greatly after the initial charge–discharge process at the cutoff of 4.3 V. SE coatings still attach on the LiCoO₂ particles (Figure 7E), and empty spaces among particles observed in panel (Figure 7B) are similar to those observed in the as-prepared electrode (Figure 7A). Figure 7C shows that voids among LiCoO₂ particles increase after the charge–discharge process at a higher cutoff voltage of 4.6 V. A LiCoO₂ particle charged to 4.6 V suffers from deterioration. Many cracks are formed in the particle, as shown in the magnified image of panel (Figure 7F). This morphological change is attributable to (1) reduction of mechanical strength and (2) excess strain at the solid–solid interfaces among LiCoO₂ particles during their large volume expansion and/or phase transition.

FIGURE 7

Figure 7. Cross-sectional SEM images of SE-coated LiCoO₂ positive electrodes (A) as-prepared, (B) after the initial charge–discharge (cutoff voltage: 4.3 V), and (C) after the initial charge–discharge (cutoff voltage: 4.6 V). Images of (D–F), respectively, depict enlarged images of (A–C).

Figure 8 shows cross-sectional SEM images of SE-coated graphite negative electrodes; Figure 8A as-prepared and Figure 8B after the initial charge–discharge (cutoff voltage: 4.3 V). Graphite particles are close-packed in a negative electrode layer. SE coatings are observed around graphite particles. They form Li⁺ ion conduction paths through the electrode layer. After the initial charge–discharge, no obvious void is observed in the negative electrode. The close-packed electrode is present. Good solid–solid interfaces are also retained between the graphite layer and an electrolyte separator layer.

FIGURE 8

Figure 8. Cross-sectional SEM images of SE-coated graphite negative electrodes (A) as-prepared and (B) after the initial charge–discharge (cutoff voltage: 4.3 V).

As described in Section “Preparation Process,” liquid-phase synthesis of sulfide electrolytes is also useful for forming favorable electrode–electrolyte solid–solid interfaces, which can be achieved by removing solvents from electrolyte solutions. An all-solid-state cell using LiCoO₂ coated with Li₃PS₄ electrolyte via an NMF solution operates as rechargeable batteries without the addition of SE and carbon-conductive additive particles to the positive electrodes (Teragawa et al., 2014a,b). However, the cell capacities are lower than those of the cells using LiCoO₂ coated with Li₃PS₄ electrolytes by PLD, because the electrolytes synthesized via the liquid-phase has lower ionic conductivities than those prepared by PLD. The LiCoO₂ particles coated with Li₆PS₅Cl electrolytes via ethanol solution with a higher conductivity than Li₃PS₄ were therefore applied to all-solid-state cells. The weight ratio of LiCoO₂/SE layer was 92.5/7.5 in SE-coated LiCoO₂ particles. An all-solid-state cell using the electrolyte-coated LiCoO₂ shows an initial discharge capacity of 45mAh g⁻¹, which is greater than that of cells using Li₃PS₄-coated LiCoO₂ (Yubuchi et al., 2015b). SE coating via electrolyte liquids is a simple and cost-effective process. Increasing conductivity of SE at lower temperatures is necessary to improve battery performance.

Sulfur-Based Nanocomposite Positive Electrodes

Sulfur is a fascinating positive electrode with high energy density because it is an abundant resource with high theoretical capacity of 1672 mAh g⁻¹ and environmental friendliness. Lithium polysulfides (Li₂S_x), which are formed during discharge (lithiation process), are readily dissolved in organic liquid electrolytes, leading to lack of a sulfur positive electrode. Dissolution of lithium polysulfides is suppressed by absorbing sulfurs in nanocarbon pores. This approach has been studied extensively (Ji and Nazar, 2010). The use of inorganic SEs fundamentally resolves the problem.

Composite sulfur electrodes consisting of S, acetylene black (AB), and Li₂S–P₂S₅ SE powders with a weight ratio of 25/25/50 were prepared using high-energy planetary ball milling to produce favorable contacts among the three components (Nagao et al., 2011). The Li-In/S cell exhibits a large reversible capacity of greater than 1500 mAh g⁻¹ of sulfur with average potential of ca. 2.1 V (vs. Li⁺/Li). The cell with sulfur electrode shows 15 times higher capacity than the cell with LiCoO₂, although the operating potential of the former cell is almost half that of the latter cell.

Application of Li₂S as a discharge product of sulfur active material offers the important benefits of high theoretical capacity of 1167 mAh g⁻¹ and versatility of negative electrode materials without lithium sources. Because of the insulative nature of Li₂S, composite electrodes, Li₂S mixed with conductive additives, such as nanocarbons and SEs, should be prepared for the use of Li₂S as an active material. Composite Li₂S electrodes were prepared by mechanical milling. A typical weight ratio of Li₂S/AB/Li₂S–P₂S₅ SE is 25/25/50. The prepared composite gives broad peaks attributable to Li₂S in X-ray diffraction patterns. An all-solid-state cell of In/Li₂S composites is charged and then discharged at 25°C. The initial reversible capacity is 800 mAh g⁻¹ at the current density of 0.064 mA cm⁻² (Nagao et al., 2012b). The cell retained 750 mAh g⁻¹ for 10 cycles. Charge–discharge reaction mechanisms were examined using high-resolution TEM observation (Nagao et al., 2015). Figure 9 shows TEM images for the Li₂S electrodes; Figure 9A before charge–discharge test, Figure 9B after the initial charge, and Figure 9C after the initial discharge. Figure 9A shows that nanoparticles of ca. 5 nm in size with different crystal orientations are distributed randomly in the matrix consisting of amorphous SE and AB. Those nanoparticles are attributable to crystalline Li₂S. Figure 9B shows that no lattice fringes because of the crystalline Li₂S are apparent, and there exists the characteristic contrast attributable to amorphous structure in the whole region after the initial charge process. As shown in Figure 9C, lattice fringes with spacing of about 3.9 Å are clearly apparent after the initial discharge, suggesting that amorphous sulfur is converted into crystalline nanoparticles during discharge reaction. Reversible transformation between crystallization and amorphization of sulfur-based active nanoparticles is responsible for the high capacity and its retention.

FIGURE 9

Figure 9. TEM images for the Li₂S electrodes (A) before charge–discharge test, (B) after the initial charge, and (C) after the initial discharge.

The utilization of Li₂S is ca. 50% in the composite electrode. To increase the utilization of Li₂S, one strategy is the fundamental enhancement of ionic conductivity in Li₂S. A partial substitution of more polarizable iodide anion with larger ionic radii for sulfide anion in Li₂S is expected to increase conductivity by introducing lithium vacancies and by increasing the lattice constant. Solid solutions in the system Li₂S-LiI are therefore prepared using mechanical milling (Hakari et al., 2015a). Only the XRD peaks attributable to Li₂S are observed; LiI peaks disappear completely in the composition range of 0 < LiI (mol%) < 20. Figure 10A presents the composition dependence of lattice constant and conductivity at 25°C for the prepared Li₂S-LiI materials. The lattice constant increases monotonically with increased LiI content, suggesting that Li₂S-based solid solutions are prepared using a mechanical milling process. Conductivity is also enhanced by increasing the LiI content. The solid solution with 20 mol% LiI has conductivity of 2.2 × 10⁻⁶ S cm⁻¹, which is two orders of magnitude higher than that for Li₂S itself without the addition of LiI. A composite positive electrode with the 80Li₂S⋅20LiI (mol%) solid solution, vapor-grown carbon fiber (VGCF), and Li₃PS₄ glass electrolyte with a weight ratio of 50/10/40 is applied to all-solid-state cells. Figure 10B shows that an all-solid-state cell (Li-In/Li₃PS₄/80Li₂S⋅20LiI) is charged and discharged at a current density of 0.13 mA cm⁻² (0.07 C) at 25°C. The cell shows a reversible capacity of 930 mAh g⁻¹ for 50 cycles. The capacity corresponds to 80% utilization of Li₂S. It is noteworthy that the enhancement of conductivity of Li₂S is effective for increasing the utilization of the active material. This strategy is beneficial for the development of all-solid-state cells with a higher energy density.

FIGURE 10

Figure 10. (A) Composition dependence of lattice constant and conductivity at 25°C for prepared Li₂S-LiI solid solutions and (B) charge–discharge curves of an all-solid-state cell (Li-In/Li₃PS₄/80Li₂S⋅20LiI).

A SE Li₃PS₄ is ball-milled with nanocarbon, such as AB. The prepared Li₃PS₄–AB materials as a mixed conductor are useful as positive electrode acting not only as a SE but also as an active material. An all-solid-state cell with the Li₃PS₄–AB composite positive electrode is charged and discharged. Its operation voltage of ca. 2.6 V vs. Li⁺/Li is somewhat higher than that of Li₂S (Hakari et al., 2015b). Redox-active electrolytes, such as CuCl₂-dissolved solution in porous carbons, are also reported to apply to supercapacitors (Mai et al., 2013). The use of the SEs as an active material in electrode layers is effective at increasing the reversible capacity per gram of the total mass of positive electrodes.

Transition Metal Sulfide Positive Electrodes

Typical MS_x, such as TiS₂, are used as active materials in all-solid-state Li and Na batteries with sulfide SEs. Decreasing the particle size of MS_x and forming a wide contact area with both SEs and conductive additives are important for increasing the MS_x utilization.

Monodispersed MS_x nanoparticles are prepared using a so-called “hot-soap” technique using high-boiling point solvents as a reaction medium. Particle morphology and size can be controlled by choosing the reaction conditions and combinations of coordinating or non-coordinating solvents. NiS particles of 50 nm size were prepared using thermal decomposition of nickel acetylacetonate in a mixed solution of 1-dodecanethiol as a sulfur source and 1-octadecene as a non-coordinating solvent at 280°C for 5 h (Aso et al., 2011). The NiS nanoparticles are crystallized directly on a carbon fiber (VGCF) by adding VGCF to a liquid medium. Good adhesion between NiS and carbon is achieved (Aso et al., 2012). Sulfide electrolyte coating on NiS–VGCF was produced using the PLD method. All-solid-state cells with the prepared NiS composite electrode operate as a secondary battery at 25°C, suggesting that electron and Li⁺ ion conduction paths are formed in the composite electrodes (Aso et al., 2013).

To increase the positive electrode capacity, sulfur-rich MS_x are desired. For example, titanium trisulfide TiS₃ shows a higher capacity than that of TiS₂, because additional sulfurs in TiS₃ contribute to the redox reaction that occurs during charge–discharge processes (Hayashi et al., 2012a). Amorphous NbS_x (x = 3, 4, 5) are prepared mechanochemically. Electrochemical cells with an organic liquid electrolyte using the amorphous NbS_x (x = 3, 4, 5) show higher discharge capacities with an increase in the sulfur content of NbS_x (Sakuda et al., 2014). Amorphous TiS₃ (a-TiS₃) retains a higher capacity than crystalline TiS₃ in all-solid-state lithium cells. The crystal structure of TiS₃ is partially deteriorated at the initial cycle, leading to an irreversible capacity in the cell with crystalline TiS₃ (Matsuyama et al., 2016a,b). It is noteworthy that sulfur-rich amorphous MS_x with electrical conductivity are promising for use as positive electrodes instead of sulfur active materials.

Figure 11A shows the first and tenth charge–discharge curves of all-solid-state lithium cell Li-In/a-TiS₃ at 0.013 mA cm⁻² at 25°C. The right side ordinate axis represents the electrode potential vs. Li⁺/Li, as calculated based on the potential difference between the Li-In and Li electrode (0.62 V). The cell with a-TiS₃ positive electrode including no carbon-conductive additives and SEs shows a reversible capacity for 10 cycles of about 510 mAh g⁻¹ of a-TiS₃, which equals the weight of the total positive electrode. The capacity corresponds to storage of about 3M Li to a-TiS₃. Electronic structural analyses using S2p XPS and S K-edge XANES reveal that a reversible sulfur redox in a-TiS₃ appears mainly during charge–discharge processes and contributes to its good capacity retention (Matsuyama et al., 2016a,b).

FIGURE 11

Figure 11. (A) Charge–discharge curves of all-solid-state lithium cell Li-In/amorphous TiS₃ and (B) charge–discharge curves of all-solid-state sodium cell Na₁₅Sn₄/amorphous TiS₃ with or without AB.

Amorphous TiS₃ is also applicable to all-solid-state sodium batteries. An all-solid-state sodium cell using a-TiS₃ shows capacity higher than 300 mAh g⁻¹ at the first discharge process, as presented in Figure 11B (Tanibata et al., 2015a). The composite positive electrode consisting of a-TiS₃ and the cubic Na₃PS₄ electrolyte with a weight ratio of 40/60 is used. Na₁₅Sn₄ alloy and the cubic Na₃PS₄ electrolyte are used, respectively, as a negative electrode and a separator layer. The cell capacity decreases gradually during discharge–charge cycles. SEM–EDX analysis reveals that the a-TiS₃ particles aggregate after the cycles, and that resistance in the a-TiS₃ composite electrode increases during cycles. To secure electron conduction paths to a-TiS₃, 6 wt% AB is added to the positive electrode. Good capacity retention is achieved in the cell using a-TiS₃ electrode with AB. The added AB particles might prevent the shutoff of electron conduction paths in the composite electrode and give a buffer space for the volume change of a-TiS₃ particles for preserving adhesion among the particles. The addition of AB to the Na₁₅Sn₄ electrode also suppresses cell resistance after charge–discharge. In addition, the replacement of Na₃PS₄ glass–ceramic by 94Na₃PS₄⋅6Na₄SiS₄ (mol%) glass–ceramic with higher conductivity decreases cell resistance and increases the rate performance of all-solid-state cells (Tanibata et al., 2015b). Reversible capacity in all-solid-state Na cells is less than that in all-solid-state Li cells, as shown in Figures 11A,B. The lower conductivity of SEs and loss of particle contacts because of the larger volume change are responsible for a smaller capacity for the Na cells. Approaches that are used to prepare composite positive electrodes suitable for insertion/de-insertion of Na⁺ ion with larger ionic radius must be developed to improve battery performance in all-solid-state Na batteries.

Interface Modification for Li Metal Negative Electrode

To achieve high energy density of all-solid-state Li batteries, the final goal is the use of lithium metal as a negative electrode. Lithium metal is an ultimate negative electrode because of a large theoretical capacity of 3861 mAh g⁻¹ and the lowest electrochemical potential of −3.04 V vs. SHE. However, a possibility of fatal problems emerged by short-circuit with dendrite formation prevent the practical use of Li metal negative electrode in lithium cells with conventional liquid and polymer electrolytes. Combination with inorganic SEs is expected to resolve the problem. In fact, thin-film solid-state batteries with Li negative electrode have excellent cycle life without capacity fading (Bates et al., 1993). The rate of utilization of Li electrodes is not high in thin-film batteries because positive electrodes such as LiCoO₂ have lithium sources, which are mainly used for a charge–discharge process.

Compared with thin-film batteries fabricated by gas-phase deposition, insufficient Li–SE interfaces in bulk-type all-solid-state batteries prepared by cold pressing are important issues that must be resolved. It has been revealed by in situ SEM observation that lithium is deposited through grain boundaries and voids in the SE (Nagao et al., 2013). Insertion of a Li-alloy thin layer at the interface between the Li electrode and SE layers brought about stable Li dissolution and deposition in the all-solid-state Li metal cells (Hiratani et al., 1988; Okita et al., 2011; Nagao et al., 2012a). These interface modifications are effective at establishing homogeneous interfaces between the Li metal and SEs.

Intensive utilization of Li is important for achieving high energy density of all-solid-state lithium metal batteries. Insertion of a Au thin film at the Li–SE interface is effective for increasing Li utilization (Kato et al., 2016a). Li and Au thin films were formed on a pelletized Li₃PS₄ glass electrolyte by vacuum evaporation. Galvanostatic cycling tests for the Li/Li₃PS₄/Li cell are presented in Figure 12A. At the initial cycle, the utilization of Li metal is about 40%, which is a higher rate of utilization of Li metal for thin-film batteries (about 20%). However, utilization of Li metal for the cell decreases rapidly after five cycles. The Li thin-film morphology is rough after galvanostatic cycling tests, as shown in the SEM image (Figure 12C), indicating that inhomogeneous Li dissolution–deposition reactions occur. The cell Li/Au/Li₃PS₄/Au/Li has about 35% Li utilization at the initial cycle and retains about 25% after the fifth cycle, as presented in Figure 12B. The morphology of Li metal after Li dissolution–deposition reaction became more uniform, as shown in the SEM image of panel (Figure 12D), compared with the cell without Au thin films. The insertion of Au film to a Li–SE interface is a first step for improving the cyclability of Li deposition–dissolution reactions with high Li utilization in all-solid-state lithium metal batteries. Intensive studies are in progress.

FIGURE 12

Figure 12. Galvanostatic cycling tests for (A) a Li symmetric Li/Li₃PS₄/Li cell and (B) a Li/Au/Li₃PS₄/Au/Li cell. SEM images of the surface of (C) a Li/Li₃PS₄/Li cell and (D) a Li/Au/Li₃PS₄/Au/Li cell after galvanostatic cycling tests.

Concluding Remarks

We have reviewed recent developments related to sulfide SEs and interface formation processes for all-solid-state rechargeable batteries. The conductivity of sulfide Li⁺ ion conductors, such as Li₁₀GeP₂S₁₂, Li₇P₃S₁₁, and Li_9.54Si_1.74P_1.44S_11.7Cl_0.3, has already reached 10⁻² S cm⁻¹ at room temperature. Conductivities of sulfide Na⁺ ion conductors are lower than those of Li⁺ ion conductors at the present stage. The highest conductivity of 10⁻³ S cm⁻¹ was obtained for Na₃PSe₄. Sulfide glass electrolytes with high alkali content were prepared by high-energy ball milling process. Metastable phases, such as Li₇P₃S₁₁ and cubic Na₃PS₄, with high conductivity were formed by careful crystallization of the prepared glasses. Sulfide electrolytes can be synthesized via liquid-phase processing, which is useful for coating application. Sulfide glass electrolytes have favorable formability, and Young’s modulus for forming good electrode–electrolyte interfaces achieving rapid charge transfer in bulk-type all-solid-state batteries. Chemical stability in air is a great shortcoming of using a sulfide electrolyte. The composition of Li₃PS₄ has superior chemical stability in the Li₂S–P₂S₅ binary system. The higher chemical stability is achieved by partial substitution of oxygen for sulfur. Selecting composition and designing structure for sulfide electrolytes is expected to improve conductivity, mechanical properties, and chemical stability further. As Na⁺ ion conductors, higher conductivity of more than 10⁻² S cm⁻¹ is predicted for Sn-substituted cubic Na₃PS₄, but the conductivity has not been achieved experimentally. Further studies seeking new electrolytes and suitable preparation processes must be undertaken.

Coating of SE and preparation of nanocomposites are useful for forming favorable solid–solid interface in an electrode layer for bulk-type all-solid-state batteries. Coating of sulfide electrolytes on LiCoO₂ or graphite particles using gas-phase or liquid-phase techniques is effective for increasing solid–solid contact area using extremely small amounts of electrolytes. Preparation of nanocomposites using high-energy ball milling is useful for sulfur or Li₂S active materials with an insulative nature. Conductivity enhancement of Li₂S by combination with LiI contributes to the improvement of Li₂S utilization. Amorphous MS_x, such as amorphous TiS₃, are attractive as mixed conductors with large capacity in all-solid-state batteries. Interface modification for Li metal negative electrode with Au thin film improves the cycle performance of Li dissolution–deposition while maintaining a high rate of utilization. For further improvement of electrochemical performance of the batteries, facile approaches achieving favorable electrode–electrolyte interfaces with large contact area will be developed. Controlling the size, morphology, and dispersibility of both electrolyte and electrode particles and selecting suitable electrolytes for maintaining close solid–solid contacts during charge–discharge processes will be assessed in future studies.

Author Contributions

AH contributed to the preparation of the manuscript. AS and MT contributed to the discussions about research results.

Conflict of Interest Statement

The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Acknowledgments

This research is supported by a Grant-in-Aid for Scientific Research from the Ministry of Education, Culture, Sports, Science and Technology (MEXT) of Japan. In particular, the research about all-solid-state Li batteries was financially supported by the Japan Science and Technology Agency (JST), Advanced Low Carbon Technology Research and Development Program (ALCA), Specially Promoted Research for Innovative Next Generation Batteries (SPRING) Project, while the research about all-solid-state Na batteries was supported by the MEXT program “Elements Strategy Initiative for Catalysts and Batteries (ESICB).”

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