Synthesis and Structures of Novel Solid-State Electrolytes

Two classes of new materials possessing ion conductivity have been developed: a lithium ion conductor and a hydride ion conductor. Conventional perovskite and ordered rock-salt structures were adopted as frameworks for lithium migration, and electrochemically stable elements such as Al, Ga, Ta, and Sc were used in the materials to facilitate their use as low-potential negative electrodes. New compositions of (Li0.25Sr0.625V(Li,Sr)0.125)(Ga0.25Ta0.75)O3, and Li0.9Sc0.9Zr0.1O2 were found to be novel oxide-based lithium ion conductors. Oxyhydrides with K2NiF4-type structures were synthesized via a high-pressure synthesis method and their use in pure hydride ion conduction was demonstrated. The La2–x–y Sr x+y LiH1–x+y O3–y oxyhydrides showed wide composition ranges of solid solution formation and the conductivity increased with anion vacancies or the introduction of interstitial hydride ions. The performance of an all-solid-state TiH2/o-La2LiHO3 (x = y = 0, o: orthorhombic)/Ti cell provided conclusive evidence of pure H– conduction.


Novel Solid-State Electrolytes
Solid materials exhibiting purely ionic conduction are used as solid-state electrolytes in a wide variety of electrochemical devices and chemical sensors, with the corresponding charge carriers being specific ions such as H + , Cu + , Ag + , Na + , Li + , F -, and O 2-. The resulting charged ion flow in electrolytes creates an electric current that drives the device, the characteristics and performance of which are thus influenced by the nature of the charge carriers. Generally, in view of their small ionic radii, cations migrate easily in solid electrolytes, showing facile diffusion. For example, silver and copper ion solid electrolytes, such as RbAg 4 I 5 and Rb 4 Cu 16 I 7 Cl 13 , show extremely high ionic conductivities of > 100 mS cm −1 at room temperature [1][2][3]. Moreover, the recently developed lithium ion conductors (Li 10 GeP 2 S 12 , LGPS) have achieved room temperature conductivities of >10 mS cm −1 [4,5], with Li-based all-solid-state batteries reported to exhibit exceptionally good power characteristics. On the other hand, newly developed materials such as hydride ion conductors have expanded the research field and the scope of available energy devices [6,7]. In this section, we focus on Li + and Has charge carriers and describe the structural characteristics of the corresponding newly developed materials.

Lithium Ion Conductors
Lithium ion conductors continue to attract much attention owing to their practical applications in all-solid-state lithium batteries [5,8]. A wide variety of such conductors exists (e.g., LISICON, perovskite, garnet, glass, glass ceramics, thio-LISICON, and LGPS), some of which were developed in the 1970s [4,[9][10][11][12][13][14]. For instance, LGPS-based materials (σ > 10 mS cm −1 at 25°C) enable high-power operation of solid-state lithium batteries; this is an intrinsic merit of solid-state systems, in addition to their safety and reliability. However, sulfide-based solids are sensitive to atmospheric moisture. As a result, most current research focuses on oxide-based materials, in order to satisfy the requirements of practical applications and engineering processes.
Novel ion conductors are typically developed using three methods: (i) element substitution-based, (ii) structure-based, and (iii) composition-based material searches. Approach (i) relies on existing materials with ionic conductivity of the target charge carrier [15], which are amenable to tuning of their physical and electrochemical properties [16]. Therefore, although it is relatively easy to find new materials using this method, remarkable performance improvements are difficult to achieve. Approach (ii) is initiated by selecting a suitable crystal structure candidate for ion diffusion [6,11], which can be complicated by the fact that the diffusion of the target ion in the selected structure has usually not been demonstrated.
Finally, approach (iii) is the most challenging, but also has the greatest potential to afford new materials with unique structures and properties. This approach starts with the selection of a suitable phase diagram [17]. Subsequently, materials corresponding to the chosen region in this diagram are synthesized and characterized; in certain cases, they exhibit unique structures and properties [4,18]. In this chapter, some examples of material searches are introduced.

Novel Lithium Ion-Conducting Perovskite Oxides [15]
Lithium ion-conducting solids are key materials for all-solid-state lithium batteries, which, compared with conventional liquid electrolyte-based lithium batteries, exhibit improved energy density, stability, safety, and reliability. Among the solid electrolytes that have been developed, the oxide-based ones are among the most promising candidates, owing to their high ionic conductivities and good chemical stabilities over a wide range of operating temperatures [11,19]. Lithium ion-conducting perovskites such as La (2/3)-x Li 3x TiO 3 (which exhibits an ionic conductivity above 10 −3 S cm −1 at room temperature) are considered to be particularly attractive [19]. However, the interfacial reduction of Ti +4 to Ti +3 during the electrochemical process or upon contact with lithium metal gives rise to undesirable electronic conduction. On this basis, novel perovskite-structured materials were examined, in typical example of an element substitution-based material search. As a result, the Li-Sr- Ta , making the Li cation more ionic in nature and, therefore, more mobile. The tolerance factor t, calculated based on the ionic radius of Sr 2+ at the A-site, equaled 0.9855, which was close to unity and indicated an ideal cubic perovskite-type structure.
The increased ionic conduction was confirmed to result from the introduction of vacancies at A-sites. Average bond valence sum (BVS) values were calculated for each site of the perovskite structure using refined structural data, with the average BVS for A-sites equaling 1.98. The larger average BVS of (Li 0. 25

M-Doped LiScO 2 (M = Zr, Nb, Ta) [21] as New Lithium Ion Conductors
No material has yet been discovered that satisfies all requirements imposed on lithium ion conductors as solid electrolytes for battery applications (i.e., high ionic conductivity at room temperature, chemical stability, electrochemical stability, thermal stability, and low cost). This clearly indicates the need for further research efforts in this direction. Herein, we focus on LiScO 2 , which has an ionic conductivity of 4 × 10 −9 S cm −1 at 573 K. Although this value is not overly high, the above material is still attractive in view of its enhanced thermodynamic stability in contact with lithium metal [22]. As shown in Fig. 13.3, LiScO 2 has a fractional cationic ordered rock-salt structure exhibiting tetragonal I4 1 /amd symmetry [23], which has the potential to partially rearrange depending on the synthesis conditions and the doped element [24]. Although element doping is an effective method of increasing the ionic conductivities of solid lithium ion conductors such as LiScO 2 , no corresponding investigations have been reported. Thus, in an attempt to improve the ionic conductivity of LiScO 2 by introducing lithium vacancies into its structure, this material was doped by M = Zr 4+ , Nb 5+ , and Ta 5+ , and the crystal structures and ionic conductivities of the thus prepared Li 1−y Sc 1−x M x O 2 were evaluated in detail.
were obtained via a solid-state reaction (sintering at 1073-1623 K for 1-12 h in air). Their impedance spectra and temperature-dependent conductivities are presented in Fig. 13.4. The  [23], with blue octahedra and green spheres indicating ScO 6 and Li, respectively conductivities were calculated from the corresponding impedance spectra, which comprised semicircles and spikes corresponding to contributions of the bulk and grain boundary and the electrode, respectively. The bulk and grain boundary contributions could not be separated and were therefore calculated together.
Resistances were calculated from the diameter of the aforementioned semicircles and used to determine conductivities. The diameters of these semicircles decreased upon doping, indicating the suitability of impedance spectroscopy to survey and evaluate ionic conductivities, with capacitance values corresponding to the observed semicircles being in the range of 10 −10 -10 −12 F. Table 13.2 summarizes the ionic conductivities and activation energies of Li 1−y Sc 1−x M x O 2 (x = 0.1) at 573 K, along with the values previously reported for LiScO 2 . All doped samples showed higher conductivities than the parent compound, owing to the formation of solid solutions upon aliovalent cation doping. Furthermore, this doping decreased the activation energies by more than 10%, indicating that the formation of lithium vacancies in the LiScO 2 lattice reduced the energy barrier of lithium diffusion.   were increased by doping with Zr 4+ , with refinement results showing that 10% Zr 4+ was doped at Sc 3+ sites in the above structure, in agreement with the ratio of utilized reactants. Concomitantly, lithium vacancies were probably formed to maintain the charge balance in LiScO 2 , since the doped Zr 4+ ion has a higher charge than Sc 3+ . These results demonstrate that the ionic conductivity of LiScO 2 was markedly improved by substitution with certain aliovalent cations, owing to the resulting lattice expansion and formation of lithium vacancies.

Development of Hydride Ion Conductors
Hydride ion conduction is particularly attractive, as His similar in size to fast ionic conduction-suitable oxide and fluoride ions, while exhibiting strong reducing properties (standard H − /H 2 redox potential = − 2.3 V), comparable to those of Mg/Mg 2+ (−2.4 V) (Fig. 13.5). Thus, hydride ion conductors may be applied in energy storage/conversion devices with high energy densities. To indicate a new direction for next-generation battery systems beyond lithium ion batteries and fuel cells, we herein focus on hydride ion conduction in solids.
Hydride ion conduction in CaH 2 was first described by Andresen et al. in 1977 [25], with similar reports on other materials following in later years [26][27][28][29][30][31]. However, experimental evidence of Hconduction was not obtained until Irvine et al. determined the transport number of BaH 2 by electromotive force measurements in 2015 [7]. Although alkaline earth metal hydrides such as BaH 2 act as pure Hconductors, they are also strong reducing agents. This complicates their use as solid electrolytes of energy devices, in which electrochemical stability to both oxidation and reduction is required. Indeed, these metal hydrides have not yet been applied to battery reactions. From the viewpoint of material design, the structural inflexibility of metal hydrides complicates the control of their lattice structure (which is required to create smooth transport pathways) and their conducting hydride ion content. Thus, little progress has been achieved in the development of Hconductors. We have considered oxyhydrides, in which hydride and oxide ions share anion sublattices, as prospective hydride conductors with flexible anion sublattices. Known oxyhydrides include A 2 BH x O 4-x (K 2 NiF 4 structure; A = La, Ce, Nd, Pr, Sr; B = Co, V, Li; 0 < x ≤ 1), Sr 3 Co 2 O 4.33 H 0.84 (Ruddlesden-Popper structure), ATiO 3-x H x (perovskite structure; A = Ba, Sr, Ca) [32][33][34][35][36][37], and [Ca 24 Al 28 O 64 ] 4+ ⋅ 4H -(mayenite structure) [38][39][40]. However, none of these materials display pure Hconductivity, since hydride ions have been reported to act as electron donors in oxide-based materials [38][39][40][41][42], donating electrons to their lattice and thus causing electron conduction accompanied by a characteristic change in hydrogen charge from Hto H + . Indeed, perovskite-and mayenite-type oxyhydrides predominantly exhibit electron conduction caused by the dissociation of hydride ions into electrons and protons [33,[38][39][40]43]. Taking this into consideration, preventing the above electron donation may be important for achieving pure Hconduction in the oxide framework structure. Herein, we attempted to synthesize a series of K 2 NiF 4 -type oxyhydrides, La 2-x-y Sr x+y LiH 1-x+y O 3-y (0 ≤ x ≤ 1, 0 ≤ y ≤ 2, 0 ≤ x + y ≤ 2) featuring cation sublattices that contain cations that are more electron-donating than Hand anion sublattices that allow flexible storage of H -, O 2-, and vacancies.

Hydride-Conducting Oxyhydrides La 2-X-Y Sr x+Y H 1-X+Y O 3-Y
Novel La 2-x-y Sr x+y LiH 1-x+y O 3-y oxyhydrides were synthesized by a high-temperature solid-state reaction in a cubic anvil cell [6] under high pressure to prevent the loss of light elements such as hydrogen, which can easily vaporize at high temperatures. The compositions and structures of La 2-y Sr y LiH 1+y O 3-y (y = 0, 1, 2) were determined by X-ray and neutron Rietveld analyses (Fig. 13.6). In

Hydride Ion Conductivity of La 2-X-Y Sr x+Y H 1-X+Y O 3-Y
The ionic conductivities of La 2-x-y Sr x+y LiH 1-x+y O 3-y were examined by impedance measurements. The Arrhenius plots of conductivities are shown in Fig. 13.9. In the case of La 2-y Sr y LiH 1+y O 3-y (x = 0), conductivity increased with increasing Hcontent, with the highest value of 3.2 × 10 −5 S cm −1 at 573 K observed for Sr 2 LiH 3 O (y = 2) (Fig. 13.9a). Thus, introduction of hydride ions into the anion sites of the K 2 NiF 4 structure improved ionic conductivity, confirming that these ions were primary charge carriers. Conduction was further facilitated by the introduction of vacancies, indicating that structural defects can affect ionic diffusion, as can be seen for La 2-x Sr x LiH 1-x O 3 (y = 0) and La 1-x Sr 1+x LiH 2-x O 2 (y = 1),  (Fig. 13.9b, c). To further identify the nature of the charge carriers, the electrical conductivity of La 0.6 Sr 1.4 LiH 1.6 O 2 (x = 0.4, y = 1.0) was evaluated by the Hebb-Wagner polarization method [45] at 480 and 590 K using an asymmetric (-) Pd/ La 0.6 Sr 1.4 LiH 1.6 O 2 /Mo (+) cell, with the total electrical conductivities (electrons + holes) at the irreversible Mo-electrolyte interface (2.9 × 10 −8 and 4.1 × 10 −7 S cm −1 , respectively) showing that La 0.6 Sr 1.4 LiH 1.6 O 2 is a purely ionic conductor (Fig. 13.10 and Table 13.7).

Development of Electrochemical Devices Based on Hydride Ion Conduction
To verify the occurrence of Hconduction in La 2-x-y Sr x+y LiH 1-x+y O 3-y , we constructed a Ti/o-La 2 LiHO 3 /TiH 2 all-solid-state cell and subjected it to galvanostatic discharge, with an electrode configuration (powdered mixture of electrode and electrolyte materials) similar to that previously used in an all-solid-state lithium battery [46]. Figure 13.11a shows the discharge curve of the cell, revealing a constant discharge current of 0.5 μA at 300°C. Moreover, the cell showed an initial open circuit voltage of 0.28 V, which was consistent with the theoretical value calculated from the standard Gibbs energy of formation of TiH 2 [47]. During the electrochemical reaction, the cell voltage rapidly dropped from 0.28 to 0.06 V and then gradually decreased to 0.0 V. The initial steep drop corresponded to an increase in hydride ion content at the anode, owing to the following constant current discharge reaction: with the cathode reaction represented as: The occurrence of these discharge reactions was confirmed by analysis of the produced phases. Figure 13.11b shows the synchrotron X-ray diffraction patterns of the cathode, electrolyte, and anode materials before and after the reaction. The absence of any variation in the diffraction patterns of the electrolyte indicates that the La 2 LiHO 3 electrolyte was stable in contact with the Ti and TiH 2 electrodes during the reaction. Conversely, phase changes were observed for the cathode and anode materials, as expected from the Ti-H phase diagram [47], where the δ-TiH 2 (Fm 3 m) phase releases hydrogen and is transformed into α-Ti (P6 3 /mmc), passing through a two-phase (α-TiH b + δ-TiH 2-a ) coexistence region found below ∼ 573 K. In the case of the cathode, additional diffraction peaks corresponding to species with P6 3 /mmc symmetry were detected, and the signals of TiH 2 shifted to a higher angle, indicating that the release of hydrogen from TiH 2 induced lattice shrinkage.
In the case of the anode, peaks corresponding to species with Fm 3̄m symmetry were detected. Thus, the results indicate that during the electrochemical reaction, hydride ions were released from the TiH 2 cathode and diffused into the Ti anode through o-La 2 LiHO 3 . The present success in the construction of an all-solid-state electrochemical cell exhibiting Hdiffusion confirms not only the ability of oxyhydrides to act as Hsolid electrolytes, but also the possibility of developing electrochemical solid devices based on Hconduction.

Ambient-Pressure Synthesis of H --Conductive Oxyhydrides
The abovementioned high-pressure method is efficient for synthesizing oxyhydrides, owing to its ability to inhibit hydrogen desorption from the starting materials during sintering. However, in order to apply Hconductors to electrochemical devices, a simple synthetic protocol needs to be established for oxyhydrides, in parallel with the development of highly H --conductive novel materials. Here, we The starting materials (which were identical to those used in the high-pressure method) were pelletized and placed in a sealed sample container made of stainless steel, with subsequent sintering performed at 650°C for 6 h under H 2 . Figure 13.12 shows the X-ray diffraction patterns of LaSrLiH 2 O 2 synthesized using different amounts of LiH (stoichiometric, 20, 50, and 100 wt% excess), with the main diffraction peaks corresponding to the space group of LaSrLiH 2 O 2 , i.e., I4/ mmm. However, small diffraction peaks indexed to SrO, SrH 2 , and/or La 2 O 3 , which were present in the raw starting materials, were observed for samples synthesized using a small excess or no excess of LiH (stoichiometric, 20 wt%, and 50 wt%). The amount of residual starting materials decreased as the amount of LiH increased, with LaSrLiH 2 O 2 obtained as a single phase only at a 100 wt% excess. In addition, excess LiH improved the crystallinity of LaSrLiH 2 O 2 , i.e., the magnification of the normalized 004 peaks (Fig. 13.12) showed that their full width at half maximum decreased as the amount of LiH increased. Therefore, excess LiH not only prevented the loss of lithium and hydrogen during the synthesis of LaSrLiH 2 O 2 , but also acted as a flux for reducing the synthesis temperature.
Crystal structure analysis revealed that the sample prepared under ambient pressure had almost the same structure as the high-pressure one, with the refined site occupancies of each atom indicating that the former exhibited a nearly stoichiometric composition without vacancies. However, mixing of H − and O in the 4c axial anion site (g(H1) = 0.9361 (5) and g(O1) = 1 − g(H1)) in Li octahedra was detected for the ambient-pressure sample, whereas this site was exclusively occupied by H − in the high-pressure sample.
The ionic conductivity of LaSrLiH 2 O 2 synthesized at ambient pressure was evaluated by AC impedance measurements. The corresponding impedance and Arrhenius plots are shown in Fig. 13.13, with the conductivity of the high-pressure LaSrLiH 2 O 2 synthesized in our previous study also plotted for comparison [6]. The impedance plot exhibited a typical form, comprising a semicircle in the high-frequency range and a spike in the low-frequency range, which corresponded to contributions of the bulk and grain boundary and the electrode, respectively. The former contribution was estimated by fitting impedance spectra using an equivalent circuit, as shown in Fig. 13.13. For the ambient-pressure sample, the activation energy of ionic conduction was calculated as 80.7 kJ mol −1 , which is nearly equal to that observed for the high-pressure sample [6]. The total conductivity (bulk + grain boundary) of the ambient-pressure sample was determined as 3.2 × 10 −6 S cm −1 at 300°C, slightly less than that of the high-pressure sample. Given the crystal structure of LaSrLiH 2 O 2 , in which tetragonal (LiH 2 ) − and (LaSrO 2 ) + layers are alternately stacked along the c-axis, the hydride ions were expected to exhibit two-dimensional diffusion in the LiH 4 plane. Hence, the movement of H − in the crystal lattice of ambient-pressure LaSrLiH 2 O 2 may have been inhibited by the presence of oxide ions in the (LiH 2 ) − layer. Thus, we successfully synthesized LaSrLiH 2 O 2 by a conventional solid-state reaction under ambient pressure [48], with a two-fold molar excess of LiH required to obtain single-phase LaSrLiH 2 O 2 . The sample synthesized at ambient pressure exhibited a crystal structure and H − conductivity similar to those observed for the high-pressure sample, implying that the method described here should increase the applicability of H − conductors as solid electrolytes.

Concluding Remarks
This chapter outlined the properties of ion conductors and material search methods, introducing Li + and H − conductors and providing examples of material search (e.g., element substitution and structure-based methods). However, broader material variability will be required to fabricate viable electrochemical devices based on solid electrolytes, necessitating the utilization of composition-based material search, which is one of the conventional material discovery methods. The approach described is significantly influenced by the experience and intuition of researchers, and it generally takes longer than element substitution and structure-based methods. However, the recent development of theoretical calculation and material informatics methods is expected to shorten the time required [49][50][51][52][53], allowing high-speed screening of prospective compositions/structures. In some cases, this approach might be misleading, since not all theoretically predicted compositions or structures can be obtained by the present synthetic techniques, as exemplified by the failure of the composition/structure-based search in the case of Li + and H − conductors. Thus, the area of materials informatics for composition-based material search is still in its infancy, but it holds promise for the future.
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