Zr- and Ce-doped Li6Y(BO3)3 electrolyte for all-solid-state lithium-ion battery

The ionic conductivity of Li6Y(BO3)3 (LYBO) was enhanced by the substitution of tetravalent ions (Zr4+ and Ce4+) for Y3+ sites through the formation of vacancies at the Li sites, an increase in compact densification, and an increase in the Li+-ion conduction pathways in the LYBO phase. As a result, the ionic conductivity of Li5.875Y0.875Zr0.1Ce0.025(BO3)3 (ZC-LYBO) reached 1.7 × 10−5 S cm−1 at 27 °C, which was about 5 orders of magnitude higher than that of undoped Li6Y(BO3)3. ZC-LYBO possessed a large electrochemical window and was thermally stable after cosintering with a LiNi1/3Mn1/3Co1/3O2 (NMC) positive electrode. These characteristics facilitated good reversible capacities in all-solid-state batteries for both NMC positive electrodes and graphite negative electrodes via a simple cosintering process.


Introduction
The steady evolution of new crystal systems with high ionic conductivities has recently heightened expectations for the realisation of all-solid-state batteries (ASSBs). 1 Indeed, the successful demonstration of a high-rate-capacity ASSB has been reported with 10 À2 S cm À1 -class Li 10 GeP 2 S 12 (LGPS)-type suldebased electrolytes. 2 In the case of oxides, the reported conductivities of z10 À3 S cm À1 in current oxide-based crystal systems (perovskites, 3 NASICON types, 4 and garnets 5 ) have encouraged the realisation of oxide-based ASSBs because they are approaching the conductivities of non-aqueous liquid electrolytes (z10 À2 S cm À1 ). Oxide-based ASSBs are expected to be applied as on-board power sources for next-generation wireless Internet-of-Things devices because of their long-life performance and thermal stability. However, the discovery of better ion-conductive oxides is not always related to the enhancement of the performance of oxide-based ASSBs. This corresponds to another issue, that is, the difficulty of impurity-free interface formation between the active materials and oxide electrolyte during cosintering. 6-10 Therefore, oxide-based ionic conductors that are both highly conductive and thermally stable with respect to the electrode materials are required for the development of electrolytes for sintered oxide-based ASSBs.
The Li 3 BO 3 -related families have been considered as suitable electrolytes for fabricating good contacts at the electrode/ electrolyte interface by cosintering. For example, Li 3 BO 3 was used as an interfacial additive between LiCoO 2 electrode particles and a garnet-type sintered electrolyte plate, and reversible charge-discharge capacities were demonstrated for the ASSB aer cosintering at 700 C, despite the low conductivity of Li 3 BO 3 (s < 10 À7 S cm À1 at 150 C). 11,12 In a later report, a solid solution of Li 2 CO 3 and Li 3 BO 3 , Li 2.2 C 0.8 B 0.2 O 3 , was described as a superior candidate for assembling the ASSB because of its better conductivity of z10 À6 S cm À1 at 25 C. 13,14 Two other Li 3 BO 3 -related electrolytes-Li 3 BO 3 -Li 2 SO 4 amorphous materials 15 and LISICON-Li 3 BO 3 amorphous materials 16 -were also attractive candidates, but the ASSBs were assembled according to their deformability due to their amorphous character rather than their thermal stability.
Among Li 3 BO 3 -related families, Lopez-Bermudez et al. proposed a solid solution of YBO 3 and Li 3 BO 3 , that is, Li 6 Y(BO 3 ) 3 (LYBO), as a new crystal-type electrolyte candidate. 17 The observed ionic conductivity was low at 1.9 Â 10 À8 S cm À1 at 50 C because Li + -ion diffusion was permitted only by thermodynamic point-defect formation in the Li 6 Y(BO 3 ) 3 . However, based on DFT calculations of the defect-formation energies and Li + -ion diffusion barriers, they suggested the structural superiority of Li + -ion diffusion in the LYBO-type structure, and mentioned the possibility of conductivity enhancement by aliovalent substitution. 17 In the present study, we prepared tetravalent ion (Zr 4+ and Ce 4+ )-doped LYBO-type materials and conrmed an ionic conductivity enhancement to z10 À5 S cm À1 at 27 C. Furthermore, the performance of oxide-based ASSBs was successfully demonstrated with the Zr-and Ce-doped Li 5.875 -Y 0.875 Zr 0.1 Ce 0.025 (BO 3 ) 3 (ZC-LYBO) electrolyte because of its thermal stability to layered rock-salt positive electrodes, such as the LiNi 1/3 Mn 1/3 Co 1/3 O 2 (NMC) positive electrode, and its large electrochemical window.
The structural changes and purities of the samples were evaluated using synchrotron X-ray diffraction (SXRD) analysis (0.5Å) performed at BL19B2, SPring-8, Sayo, Japan (2019B1881). Rietveld crystal structure renement of the obtained SXRD patterns was performed using the RIETAN-FP soware program, 18 where a modied split pseudo-Voigt function 19 was selected for tting to best represent the prole parameters of the samples. To estimate the Li + -ion diffusion pathways in the structure, the rened structural parameters were used as the input le for bond valence site energy (BVSE) calculations, which were performed using SoBV soware. 20,21 The structural images and maps were drawn using the VESTA soware program. 22 For ion conductivity measurements, both sides of the tetravalent ion-containing LYBO compacts were polished and coated with gold (via sputtering) as blocking electrodes. The AC impedance prole was collected with a frequency response analyser (Solartron, 1296) over a frequency range of 0.1 Hz to 1 MHz at 27 C.
To measure the charge-discharge performance of the ASSB, a composite electrode powder consisting of 50 wt% active material and 50 wt% ZC-LYBO electrolyte was rst prepared by mixing in an agate mortar. The active material was either LiNi 1/ 3 Mn 1/3 Co 1/3 O 2 (NMC, Toda Kogyo Corp.) or graphite (JFE Chemical Corp.). The composite electrode powder (10 mg) and a Au plate as a current collector were placed on a ZC-LYBOelectrolyte separator of 30 mg in a 10 mm-diameter carbon die, which was heated at 550 C with an electric current under a pressure of 30 MPa in the SPS process. Lithium foil was used as the counter electrode. A polyethylene oxide (PEO)-based polymer electrolyte lm (Osaka Soda, LiTFSA/EO ¼ 0.06) was placed between the lithium foil and the electrolyte side of the composite electrode/electrolyte pellet to reduce the interfacial resistance via adhesion. 13 To conrm the reactivity of the composite electrode aer cosintering, laboratory XRD patterns were collected on a diffractometer using Cu Ka radiation (Miniex600, Rigaku). All XRD analyses were performed in the Bragg-Brentano geometry mode. The microstructure of the composite electrode was observed using eld-emission scanning electron microscopy (FE-SEM; JEOL, JSM-5500LV).
To evaluate the electrochemical stability of the ZC-LYBO electrolyte, a metal substrate (Au or Cu) was placed on a ZC-LYBO electrolyte (100 mg) in a 10 mm-diameter carbon die, which was heated at 550 C with an electric current under a pressure of 30 MPa in the SPS process. Lithium foil and a PEObased polymer electrolyte lm were used as the counter electrode and interfacial connector, respectively.

Results and discussion
Structural changes and purities of tetravalent ion-doped LYBO Fig. 1 shows the SXRD patterns of the Zr 4+ ion-containing LYBO samples, Li 6Àx Y 1Àx Zr x (BO 3 ) 3 (0 # x # 0.6). The I(1 1 À2) diffraction peak is clearly shied to a higher angle with an increasing amount of Zr 4+ until x z 0.1 in Li 6Àx Y 1Àx Zr x (BO 3 ) 3 because of the difference in ionic radii: 1.019Å for Y 3+ and 0.84 A for Zr 4+ in dodecahedral coordination. 23 In addition, the I(1 1 À2) diffraction peak is broadened with an increasing amount of Zr 4+ ions, which corresponds to the lattice distortion of the LYBO-type structure in addition to the aforementioned mismatch of ionic radii, which will be discussed later. Furthermore, the ZrO 2 impurity phase begins to appear from x z 0.1. Rietveld renements of the SXRD patterns were performed to clarify the structural changes caused by the Zr 4+ ions. The rened patterns for typical samples and the estimated parameters for all the samples are summarised in Fig. S2(a-c), and Table S1, † respectively. Fig. 2(a) shows the changes in the cell volume and full-width-at-half-maximum (FWHM) parameter U for the LYBO-type phases estimated from the renements. Herein, U is one of the simulated components for representing the FWHMs of the total diffraction patterns based on the Caglioti formula 24 under Rietveld renement.
where parameters V and W mainly depend on the diffractometer characteristics, which were estimated from the renements of a CeO 2 reference in advance, and q Bragg is the diffraction peak angle.
Thus, U represents the FWHM corresponding to the structural distortion in a sample. It should be noted that the structural distortion is difficult to discuss quantitatively in this study because a modied split pseudo-Voigt function 19 was used to rene the proles. Despite the difficulties in this case, the denite change in U due to the Zr 4+ ion is strongly related to the structural distortion until x z 0.1 in Li 6Àx Y 1Àx Zr x (BO 3 ) 3 , which is due to the substitution of smaller Zr 4+ ions at the Y 3+ sites. The shrinkage of the cell volume until x z 0.1 is also proof of Zr 4+ -ion doping. When x > 0.1, these changes are suppressed because of the limits of Zr 4+ -ion doping into the LYBO-type structure; thereaer, the ZrO 2 impurity (and Li-B-O impurities, mainly Li 6 B 4 O 9 ) is observed. Moreover, the lattice parameter b and area of the ac plane are shown separately in Fig. 2(b). The LYBO-type structure has a layered structure that is constructed by the stacking of the Li 2 Y(BO 3 ) 3 layer and the Li layer along the b-axis, as shown in Fig. 3(a). The shrinkage of the YO 8 dodecahedra upon doping with smaller Zr 4+ ions until x z 0.1 directly reduces the area of the ac plane. On the other hand, the lattice parameter b is expanded until x z 0.1 because of the decrease in Coulomb repulsion force at the Li layer due to the lack of Li + ions under the doping of Zr 4+ ions into Li 6Àx Y 1Àx -Zr x (BO 3 ) 3 . This tendency is similar to the change in the lattice parameter c during the electrochemical de-lithiation of layered rock-salt Li 1Àx CoO 2 . 25 Therefore, Li + ions are likely to be removed from the Li layer rather than from the Li 2 Y(/Zr)(BO 3 ) 3 layer by doping Zr 4+ ions into Li 6Àx Y 1Àx Zr x (BO 3 ) 3 , which is supported by previous DFT calculation results in which the Li defect energies at the Li layer are lower than those at the Li-ion conductive pathways The Li + -ion diffusion pathways in the LYBO-type structure were examined from bond-valence site energy (BVSE) maps; the ab and ac planes are shown in Fig. 3(b-d) and S3(b-d), † respectively. The BVSE maps present the theoretical mobile Li + -ion pathways as clouds with low bond valence site energies. Although the mobile Li + ions form a 2D diffusion pathway network at the Li layer on the ac plane [ Fig. S3(b) †], it is difficult for Li + ions to diffuse along the b axis via Li3 or Li5 sites in undoped Li 6 Y(BO 3 ) 3 , as shown by the red dashed boxes in Fig. 3(b). This was previously suggested from DFT calculations. 17 As shown in Fig. 3(c, d), the expansion of the diffusion pathways along the b-axis could be induced by doping Zr 4+ ions into Li 6Àx Y 1Àx Zr x (BO 3 ) 3 , which indicates that the structural  distortion in the Li 2 Y(/Zr)(BO 3 ) 3 layer contributed to expanding the diffusion pathways. The 2D diffusion pathways in the Li layer also become thin, as can be conrmed from Fig. S3(c, d). † These results indicate that Li + ions are more easily diffused by the structural distortion associated with doping Zr 4+ ions into Li 6Àx Y 1Àx Zr x (BO 3 ) 3 . The activation energies for the 2D and 3D pathways estimated from BVSE decrease from 0.57 and 0.69 eV to 0.49 and 0.59 eV, respectively, even by doping with only x ¼ 0.025 Zr 4+ ions into Li 6Àx Y 1Àx Zr x (BO 3 ) 3 . Fig. 4 shows the ionic conductivity dependence when Zr 4+ ions are incorporated into Li 6Àx Y 1Àx Zr x (BO 3 ) 3 (0 # x # 0.1). The observed conductivity at 27 C is drastically enhanced from 5.6 Â 10 À11 to 5.8 Â 10 À6 S cm À1 by doping with x ¼ 0.025 Zr 4+ ions into Li 6Àx Y 1Àx Zr x (BO 3 ) 3 . Li + -ion diffusion in Li-ion fully occupied Li 6 Y(BO 3 ) 3 occurs only through the formation of a few thermodynamic point defects. In contrast, the substitution of tetravalent Zr 4+ ions in the Y 3+ sites results in the intentional formation of vacancies at Li sites in the structure, which contributes to the conductivity enhancement. Additionally, the observed conductivities are sensitive to the relative densities of the Li 6Àx Y 1Àx -Zr x (BO 3 ) 3 compacts, which are summarised in Table S2. † For example, the observed conductivity in a low-density Li 3.5 Ge 0.75 -S 0.25 O 4 compact (79%) was over a hundred times lower than a high density compact ($90%). 26 Therefore, the increase in relative density upon doping with Zr 4+ ions, from 77% (x ¼ 0) to $90% (x > 0), also contributes to the ionic conductivity enhancement. Moreover, the expansion of 3D diffusion pathways by Zr 4+ doping, as estimated by the Rietveld analyses, would also affect the enhancement. To conrm this effect, the structure and conductivity of Ce-doped Li 5.975 Y 0.975 Ce 0.025 (BO 3 ) 3 were also examined. The ionic radius for the Ce 4+ ion dopant in dodecahedral coordination is 0.97Å, which is more comparable to that for the Y 3+ ion (1.019Å) than the Zr 4+ ion dopant (0.84Å). 23 Therefore, the lattice distortion by doping the tetravalent ion is supressed; the lattice distortion in Ce-doped Li 5.975 Y 0.975 -Ce 0.025 (BO 3 ) 3 (U ¼ 0.02354 deg 2 ) is closer to that in undoped Li 6 Y(BO 3 ) 3 (U ¼ 0.01488 deg 2 ) than that in Zr-doped Li 5.975 -Y 0.975 Zr 0.025 (BO 3 ) 3 (U ¼ 0.119 deg 2 ), as can be conrmed in Table  S1. † Considering the conductivities shown in Table S3, † that of the LYBO-type structure is drastically enhanced by 0.025 Ce 4+ doping (Li 5.975 Y 0.975 Ce 0.025 (BO 3 ) 3 ; s ¼ 6.9 Â 10 À7 S cm À1 at 27 C), which is due to the formation of vacancies at Li sites in the structure and the increase in compact densication, as discussed above. However, the conductivity of the Ce 4+ -ion doped sample is over ten times lower than that of the Zr 4+ -ion doped one. This difference indicates that the structural distortion associated with the doping of the larger Zr 4+ ions certainly affects the enhancement of the conductivity. The differences in the Li + -ion diffusion pathways estimated from the BVSE maps also support the aforementioned results: for the less-structurally-distorted Ce 4+ion doped sample ( Fig. S4(a, b) †), the diffusion pathway along the b-axis in the BVSE map is nearly unchanged. Therefore, Zr 4+ -ion doping exerts three types of effects that enhance the conductivity of the LYBO-type structure: (1) the formation of vacancies at Li sites, (2) an increase in compact densication, and (3) an increase in the Li + -ion conduction pathways in the LYBO phase associated with structural distortion. The highest conductivity in Zr 4+ -iondoped Li 6Àx Y 1Àx Zr x (BO 3 ) 3 at 27 C is 1.4 Â 10 À5 S cm À1 at x ¼ 0.10. However, Zr 4+ -ion doping achieves a limit at x z 0.1 in Li 6Àx Y 1Àx Zr x (BO 3 ) 3 , which would be due to the large mismatch of ionic radii. Therefore, the conductivity decreases for x > 0.1 (Table S3 † The conductivity of Zr,Ce-doped Li 5.875 Y 0.875 Zr 0.1 Ce 0.025 Zr 0.1 (-BO 3 ) 3 (ZC-LYBO) was also measured, and the obtained value of 1.7 Â 10 À5 S cm À1 at 27 C is the highest conductivity of LYBOtype oxides to date. The structural information estimated from Rietveld analysis, Li + -ion diffusion pathways estimated from BVSE maps, and relative densities of the compacts are summarised in Table S1, Fig. S3(d), and Table S2. † Compared with Zrdoped Li 5.9 Y 0.9 Zr 0.1 (BO 3 ) 3 , a slight enhancement in conductivity is conrmed, but the reason is difficult to determine in the present study. The conductivity measurement results indicate the possibility of further enhancement by other aliovalent substitutions. Moreover, the control of structural distortion is one of the keys for enhancing the conductivity of the LYBO-type structure. It should be noted that the activation energies estimated from Arrhenius plots (Fig. S5 †) were 0.43 eV for both Zr-doped Li 5.9 -Y 0.9 Zr 0.1 (BO 3 ) 3 and Zr,Ce-doped Li 5.875 Y 0.875 Ce 0.025 Zr 0.1 (BO 3 ) 3 (ZC-LYBO). Based on its high conductivity, ZC-LYBO was used as the electrolyte in subsequent ASSB studies.

Ionic conductivities
Thermal stability aer cosintering with layered rock-salt oxide active material The reactivity of the layered rock-salt LiNi 1/3 Mn 1/3 Co 1/3 O 2 (NMC) electrode and ZC-LYBO electrolyte aer cosintering was evaluated by laboratory XRD patterns (Fig. 5(a)). None of the XRD peaks are associated with impurities, and no signicant peak shis aer co-sintering are observed for either the NMC electrode or ZC-LYBO electrolyte, as shown for a typical peak in Fig. 5(b). Thus, interdiffusion between the NMC and ZC-LYBO was minimal, with no penetration of the bulk materials during cosintering. This result indicates that ZC-LYBO is thermally stable with a layered rock-salt NMC electrode.
The cross-sectional microstructure of the NMC + ZC-LYBO composite electrode aer the 550 C SPS process was also investigated. As shown in Fig. 6(a), the NMC electrode particles (light grey) are embedded in the ZC-LYBO electrolyte (dark grey). Moreover, no micropores are present near the NMC electrodes because the ZC-LYBO electrolyte was densied aer SPS. No evidence of diffusion is found in the EDX mappings for Co [NMC electrode, Fig. 6(b)], Y, and Zr [ZC-LYBO electrolyte, Fig. 6(c and d)]. This indicates that the ZC-LYBO electrolyte and NMC electrode can be co-sintered without the formation of any interfacial impurities, which is one of the rare characteristics of the electrolyte and is not possible with other popular oxideelectrolyte candidates (e.g. perovskite-type conductors, NASI-CONs, and garnet-type conductors). Fig. 7(a) shows the charge-discharge proles for an ASSB (NMC + ZC-LYBO composite electrodejZC-LYBO separatorjdry polymerjLi metal) assembled via SPS. A reversible capacity of $120 mA h g À1 is observed at 60 C owing to the impurity-free sintered interface between the NMC and ZC-LYBO. The chargedischarge capacity retentions at various current densities are shown in Fig. 7(b). Although the capacity decreases at higher current densities, relatively stable capacity retention under cycling for each current density is observed for the ASSB. In fact, the rate performance of the ASSB using the ZC-LYBO electrolyte is lower than that using LISICON-type Li 3.5 Ge 0.5 V 0.5 O 4 (LGVO), which has also been reported as a thermally stable electrolyte with an NMC electrode. 27 This is due to the lower conductivity of ZC-LYBO compared to Li 3.5 Ge 0.5 V 0.5 O 4 (i.e. 9.6 Â 10 À5 S cm À1 at   25 C). 27 However, a broader acceptance of aliovalent substitutions in the less-reported LYBO-type electrolytes could lead to the development of a more highly ion-conductive LYBO family, analogous to the advances achieved with other electrolyte candidates. 1 Electrochemical window Fig. 8(a) shows the cyclic voltammograms (CVs) of test cells (metal substratejZC-LYBO electrolytejdry polymerjLi metal) assembled by SPS. We evaluated the electrochemical stability of the ZC-LYBO electrolyte in the low-voltage range from 3.0 to À0.5 V vs. Li/Li + and in the high-voltage range from 2.0 to 5.0 V vs. Li/Li + with a Cu substrate (red) and an Au substrate (blue), respectively. The ZC-LYBO electrolyte is stable at high voltages. Au is kinetically stable within a wide electrochemical window but produces Li x Au y alloys at voltages below 0.2 V vs. Li/Li + when in contact with Li metal. 28,29 Therefore, Cu was used as the substrate to evaluate the low-voltage stability. 30 Cathodic and anodic currents corresponding to Li metal deposition (Li + + e À / Li) and dissolution (Li / Li + + e À ), respectively, are observed at 0 V vs. Li/Li + . Although the ZC-LYBO electrolyte is nearly electrochemically stable in the range from 0.0 to 5.0 V vs. Li/Li + , a slight cathodic peak can be observed below z1.0 V vs. Li/Li + .
Thus, the CVs under 5 cycles in the range from 3.0 to 0.0 V vs. Li/ Li + were also recorded ( Fig. 8(b), insert). A slight cathodic peak continuously appears in each CV cycle and is irreversible, which would result in decomposition on the surface of the ZC-LYBO electrolyte at low voltage. Fig. 8(b) shows the conductivity change for the LYBO electrolyte aer each CV cycle. Despite the continuous appearance of a cathodic peak, the Nyquist plots, including the resistance in the ZC-LYBO electrolyte, change less. Therefore, ZC-LYBO can also be used as an electrolyte with a low-voltage negative electrode.
The charge-discharge proles for a battery (Cujgraphite negative electrode + ZC-LYBO composite electrodejZC-LYBO separatorjdry polymerjLi metal) assembled by the SPS process are shown in Fig. 9. A typical reversible Li-ion (de)intercalation reaction for a graphite negative electrode is observed, and a rst discharge capacity of $200 mA h g À1 was obtained along with a large irreversible capacity for each cycle. The irreversible capacity is related to the slight anodic reaction below 1.0 V Li/ Li + (Fig. 8).
Among the various oxide-type electrolytes, ZC-LYBO is rare; it is possible to assemble ASSBs by sintering with both layered lock-salt positive electrodes such as NMC and low-voltage negative electrodes such as graphite. LISICON-type Li 3.5 Ge 0.5 -V 0.5 O 4 (LGVO) can also be cosintered with layered rock-salt positive electrodes but decompose at low voltage because of the presence of Ge 4+ and V 5+ ions. 27 In fact, other Li 3 BO 3 -related families such as Li 2.2 C 0.8 B 0.2 O 3 (ref. 13 and 14) and LISICON-Li 3 BO 3 (ref. 16) are stable at low voltage, although conductivities are limited to z10 À6 S cm À1 at 25 C. Therefore, the LYBO-type crystal system may be an attractive candidate for assembling ASSBs by sintering aer further enhancing the conductivity via aliovalent substitution. Even the ZC-LYBO in the present study can be used as an ion-conductive ceramic binder for sintering ASSBs instead of low-ion-conductive Li 3 BO 3 , 12 which has been used to connect LiCoO 2 electrodes and high-ion-conductive garnet-type electrolytes. 11

Conclusions
We substituted tetravalent ions (Zr 4+ and Ce 4+ ) for the Y 3+ sites in the Li 6 Y(BO 3 ) 3 (LYBO)-type structure to enhance conductivity. Fig. 8 (a) Cyclic voltammograms (CVs) of test cells (metal sub-stratejZC-LYBO electrolytejdry polymerjLi metal) prepared by the 550 C-SPS process at a scan rate of 100 mV s À1 at 60 C. Cu and Au metal substrates were used for testing the cathodic and the anodic reactions, respectively. (b) Nyquist plots of the test cell using Cu substrate at 60 C after the CV cycling and the CV profile at each cycle. The ionic conductivities of Zr-doped Li 5.9 Y 0.9 Zr 0.1 (BO 3 ) 3 and Zr,Ce-doped Li 5.875 Y 0.875 Ce 0.025 Zr 0.1 (BO 3 ) 3 (ZC-LYBO) were respectively 1.4 Â 10 À5 and 1.7 Â 10 À5 S cm À1 at 27 C, which are some of highest conductivities reported for Li 3 BO 3 -related electrolyte candidates. The effects of Zr 4+ -ion doping on the conductivity of the LYBO-type structure were revealed from the structural information estimated from Rietveld analysis, relative density of compacts, and Li + -ion diffusion pathways estimated from BVSE maps, and included: (1) the formation of vacancies at Li sites, (2) the increase of compact densication, and (3) an increase in the Li + -ion conduction pathways in the LYBO phase associated with structural distortion.
Furthermore, sintered ASSBs using ZC-LYBO as an electrolyte successfully performed with both a LiNi 1/3 Mn 1/3 Co 1/3 O 2 (NMC) positive electrode and a graphite negative electrode. This was due to the thermal stability of the layered rock-salt oxide and the electrochemical stability of the ZC-LYBO at low voltages. The conductivity of the LYBO-type electrolyte could be further enhanced by other aliovalent substitutions for the practical use of sintered ASSBs. The control of the structural distortion associated with the dopant size, which was revealed in this study, could play an important role in this enhancement.

Conflicts of interest
There are no conicts to declare.