# Fileset [s41467-025-62860-1.pdf](https://mdr.nims.go.jp/filesets/15f5f38b-4fe0-4ff3-85a6-3ce3552a009c/download) ## Creator [Daisuke Ito](https://orcid.org/0009-0006-9296-1926), [Naoaki Kuwata](https://orcid.org/0000-0002-0736-6967), Seiji Takemoto, [Kazuhiro Kamiguchi](https://orcid.org/0009-0000-8416-6590), [Gen Hasegawa](https://orcid.org/0000-0002-9297-6902), [Kazunori Takada](https://orcid.org/0000-0001-7568-1806) ## Rights [Creative Commons BY-NC-ND Attribution-NonCommercial-NoDerivs 4.0 International](https://creativecommons.org/licenses/by-nc-nd/4.0/) ## Other metadata [Lattice-matched antiperovskite-perovskite system toward all-solid-state batteries](https://mdr.nims.go.jp/datasets/0966fe59-ffda-4093-9bae-69a1e1248f07) ## Fulltext Lattice-matched antiperovskite-perovskite system toward all-solid-state batteriesArticle https://doi.org/10.1038/s41467-025-62860-1Lattice-matched antiperovskite-perovskitesystem toward all-solid-state batteriesDaisuke Ito 1 , Naoaki Kuwata 2, Seiji Takemoto1, Kazuhiro Kamiguchi 1,Gen Hasegawa 2 & Kazunori Takada 2Inorganic solid electrolytes have emerged as promising candidates for realiz-ing all-solid-state batteries because they eliminate flammable, low boiling-point liquids in lithium-ion battery cells, improving safety and cycle life. In thisstudy, we present a highly lattice-matched composite solid electrolyte con-sisting of an antiperovskite-perovskite system, offering the benefits of bothantiperovskites as melt-infiltratable solid electrolytes and perovskites as fast-ion conductors. Atomistic simulations predict significant lithium-ion diffusionat the interface between cubic Li2OHCl and Li0.31La0.56TiO3. The incorporationof fluorine enables room-temperature operation by stabilizing the high-temperature cubic phase of Li2OHCl1-xFx and reduces the lattice mismatchratio to 0.8% at the interface through lattice contractions. The composite solidelectrolyte was synthesized via pressure-assisted melt infiltration. The solidelectrolyte effectively infiltrates conventional lithium-ion battery electrodeswhile maintaining a stable interface structure. Electrochemical testingdemonstrates promising charge-discharge characteristics, including longcycle life and rate performance. Intricate infiltration of the solid electrolyteinto an electrode structure composed of activematerials withmicrocracks andhigh surface area enables stable operation by mitigating degradation phe-nomena typically observed in liquid electrolyte-based lithium-ion batteries.High-performance energy storage demand has risen owing to thewidespread use of portable devices, electric vehicles, and renew-able energy technologies1–3. Lithium-ion batteries (LIBs) are favoredfor their energy density and long cycle life4,5. However, their safetyissues, capacity limitations, and environmental impact haveprompted the exploration of alternative technologies1,6. All-solid-state batteries (ASSBs) are a potential next-generation solution tothese concerns7,8, but technical challenges hinder their commer-cialization. They require low-temperature processing, especiallywhen utilizing oxide solid electrolytes8,9. Conventional manu-facturing involves high-temperature sintering, which hampersscalability and is incompatible with temperature-sensitive compo-nents. Using low-temperature, formable, sulfide-based solid elec-trolytes carries safety risks, as they react quickly with moisture8,generating toxic hydrogen sulfide and compromising batteryreliability. Additionally, introducing solid electrolytes into theinterior of cost-effective, environmentally friendly conventional LIBelectrodes in ASSBs remains challenging. Therefore, the search forsolid electrolytes that can be incorporated into conventional LIBelectrodes while being less toxic and not reliant on rare elements iscrucial for ASSB development.Molten-salt solid electrolytes with amelt infiltration process usingLi2OHCl antiperovskite10materials, composedof rare-element-free andenvironmentally friendly rawmaterials, have been proposed byXiao etal.11. These antiperovskite materials exhibit low melting points(523–573 K) and high wettabilities toward oxides and metals, facilitat-ing penetration into electrode voids. However, Li2OHCl undergoes astructural phase transition from a low-temperature orthorhombicReceived: 25 July 2024Accepted: 4 August 2025Check for updates1Murata Manufacturing Co., Ltd., Nagaokakyo-shi, Kyoto, Japan. 2Research Center for Energy and Environmental Materials, National Institute for MaterialsScience (NIMS), Tsukuba, Ibaraki, Japan. e-mail: daisuke.ito009@murata.comNature Communications | (2025) 16:7372 11234567890():,;1234567890():,;http://orcid.org/0009-0006-9296-1926http://orcid.org/0009-0006-9296-1926http://orcid.org/0009-0006-9296-1926http://orcid.org/0009-0006-9296-1926http://orcid.org/0009-0006-9296-1926http://orcid.org/0000-0002-0736-6967http://orcid.org/0000-0002-0736-6967http://orcid.org/0000-0002-0736-6967http://orcid.org/0000-0002-0736-6967http://orcid.org/0000-0002-0736-6967http://orcid.org/0009-0000-8416-6590http://orcid.org/0009-0000-8416-6590http://orcid.org/0009-0000-8416-6590http://orcid.org/0009-0000-8416-6590http://orcid.org/0009-0000-8416-6590http://orcid.org/0000-0002-9297-6902http://orcid.org/0000-0002-9297-6902http://orcid.org/0000-0002-9297-6902http://orcid.org/0000-0002-9297-6902http://orcid.org/0000-0002-9297-6902http://orcid.org/0000-0001-7568-1806http://orcid.org/0000-0001-7568-1806http://orcid.org/0000-0001-7568-1806http://orcid.org/0000-0001-7568-1806http://orcid.org/0000-0001-7568-1806http://crossmark.crossref.org/dialog/?doi=10.1038/s41467-025-62860-1&domain=pdfhttp://crossmark.crossref.org/dialog/?doi=10.1038/s41467-025-62860-1&domain=pdfhttp://crossmark.crossref.org/dialog/?doi=10.1038/s41467-025-62860-1&domain=pdfhttp://crossmark.crossref.org/dialog/?doi=10.1038/s41467-025-62860-1&domain=pdfmailto:daisuke.ito009@murata.comwww.nature.com/naturecommunicationsphase (o-Li2OHCl) to a high-temperature cubic phase (c-Li2OHCl) atapproximately 345 K. The o-Li2OHCl has low ionic conductivity(approximately 1 × 10-7 S cm-1)12,13, rendering room-temperature ASSBoperation problematic.Here, we propose a composite solid electrolyte utilizing anantiperovskite-type fluorine-doped c-Li2OHCl1-xFx combined with aperovskite-type Li0.31La0.56TiO3 (LLTO) with a high lattice-matchedsystem. Fluorine-doped Li2OHCl1-xFx maintains its high-temperaturephase and exhibits high ionic conductivity even at roomtemperature14,15. LLTO is a perovskite-type oxide solid electrolytecomposed of two perovskite units stacked together in a unit cell8,16.Its bulk ionic conductivity exceeds 1.0 × 10-3 S cm-1 at roomtemperature16,17, but it requires sintering above 1400K and is thereforeunsuitable as a standalone ASSB solid electrolyte17,18. A lattice-matchedcomposite solid electrolyte system can effectively utilize both Li2OHCland LLTO as solid electrolytes.In this paper, we discuss the possibility of composite solid elec-trolytes based on the lattice-matched system of Li2OHCl1-xFx/LLTOinterfaceanddemonstrate ASSBusing the composite solid electrolyteswith melt infiltration processes.ResultsModeling of lattice-matched composite solid electrolytesFigure 1a shows structure models of antiperovskite c-Li2OHCl,perovskite LLTO, garnet Li7La3Zr2O12(LLZO)8, and NASICONLi1.3Al0.3Ti1.7(PO4)3 (LATP)8. The corresponding structures are availablein Supplementary Data 1. The antiperovskite structure of c-Li2OHClexhibits similarities to the LLTO perovskite structure, which is char-acterized by a reversed arrangement of cations and anions. However,LLZO and LATP have different atomic arrangements compared to thec-Li2OHCl, implying more complex interface states with c-Li2OHCl.Reciprocal lattice projections for each crystal structure are shown inFig. 1b. Antiperovskite c-Li2OHCl (a = b = c = 3.911 Å) exhibits remark-ably high lattice matching with perovskite LLTO (a = b = 3.869Å andc = 7.742 Å (3.871 Å in the cubic unit)), achieving a mismatch ratio ofonly 1.1% for energetically favorable surfaces, indicating suitable epi-taxial relationships as a composite solid electrolyte. The detailedmismatch ratios based on energetically favorable surfaces and latticeparameters are shown in Supplementary Fig. 1.Next, we discuss the interface stability of composite solid elec-trolytes based on atomistic simulations. The detailed structural00.0050.010.0150.020.0250.03F Cl Br IInterfacial energy (eV/Å2 )-3.3-3.25-3.2-3.15-3.1-3.055657585960616263F Cl Br IFormation energy (eV/atom)Cell volume (Å3 )abc-Li2OHCl LLTO LLZO LATPc-Li2OHCl c-Li2OHCl / LLTO c-Li2OHCl / LLZO c-Li2OHCl / LATP110 100 1-10010 0-10-110 -100 -1-101-20210 200 2-10 2-20-210 -200 -2-10 -2-20-1-200-20mismatch: 1.1% mismatch: 9.5%(high-order)mismatch: 6.9%(high-order)cLattice matchingd e fOOOOOOF lClCLiLaTi00.020.040.060.080.10.120.140.160.180.2Interfacial energy (eV/Å2 )c-Li2OHCl/ LLTOc-Li2OHCl/ LLZOc-Li2OHCl/ LATPContraction220120020-120-220Coherenceof atomicarrangementOZrLaLi TiAlLi OPLi HO Cl TiLaLi OFig. 1 | Atomisticmodeling of Lattice-matched composite electrolytes betweenhalogen-substituted c-Li2OHCl and oxide solid electrolytes. a Crystal structuremodels of antiperovskite c-Li2OHCl, perovskite LLTO, garnet LLZO, and NASICONLATP. Two-color spheres indicate mixed occupancy at the same crystallographicsite. The size of each color segment corresponds to the occupancy ratio.bCalculated two-dimensional reciprocal lattice projections for c-Li2OHCl, Li2OHCl/LLTO, Li2OHCl/LLZO, and Li2OHCl/LATP. White patterned dots and numericalindices in each reciprocal lattice projection represent theirMiller indices relative tothe energetically favorable (100) surface of Li2OHCl. Red, light green, and light bluedots represent projection patterns against the LLTO(001), LLZO(100), andLATP(012) surfaces, respectively. Each projection pattern is overlaid on the whitepattern of c-Li2OHCl(100). c Calculated interfacial energies between LLTO(100),LLZO(100), and LATP(012) against c-Li2OHCl(100). d Schematic illustration of theenergetically favorable interface between c-Li2OHCl(100) and LLTO(100), showingthe effects of fluorine substitution to enhance the lattice matching and coherenceof atomic arrangement. e Dependence of halogen substitution on cell volume andformation energy of c-Li2OHCl0.875X0.125 (X: F, Cl, Br, and I). f Dependence ofhalogen substitution on the interfacial energy between c-Li2OHCl(100) andLLTO(100). Source data are provided as a Source Data file.Article https://doi.org/10.1038/s41467-025-62860-1Nature Communications | (2025) 16:7372 2www.nature.com/naturecommunicationsmodeling procedure is described in the Methods section. Briefly, formodeling composite solid electrolytes, each material was optimizedusing the replica exchange Monte Carlo (REMC) method19 with thePreferred Potential (PFP)20, a machine learning potential provided byMatlantis20–22. The REMCmethod and subsequent structural relaxationfor interface modeling, along with energy evaluation during MDsimulations, were comprehensively verified by DFT calculations (Sup-plementary Fig. 2c, d, and e). The simulation results obtained in thisstudy, based on PFP, enabled a quantitative evaluation of interfacestability and the effects of Li-ion diffusion.Figure 1c shows interfacial energies between LLTO(100),LLZO(100), and LATP(012) against c-Li2OHCl(100), using energeticallyfavorable surfaces. The corresponding structures are available inSupplementary Data 2. Among these combinations, Li2OHCl(100)/LLTO(100) has the lowest interfacial energy, suggesting a thermo-dynamically stable interface. In contrast, LATP exhibits much higherinterfacial energy compared to LLTO, which is likely attributed to itspolyanionic framework containing PO4 groups that differ from theatomic arrangement of c-Li2OHCl. Figure 1d shows a schematicrepresentation of the energetically favorable interface betweenc-Li2OHCl(100) and LLTO(100). The coherence between the latticematching and atomic arrangements of these two materials minimizesatomic disorder at the interface, emphasizing their critical role informing stable interfaces for composite solid electrolytes.To further enhance lattice matching between these materials, wepropose a strategy involving halogen substitution of c-Li2OHCl withfluorine. Figure 1e shows the elemental dependence of halogen sub-stitution on the cell volume and the formation energy forc-Li2OHCl0.875X0.125 (X: F, Cl, Br, I). Substituting chlorine (ionic radius:1.81 Å) with fluorine (ionic radius: 1.33 Å) induces lattice contraction inc-Li2OHCl. The lattice parameter decreases to a = b = c = 3.87Å for Fsubstitution, reducing mismatch with LLTO to 0.1%. Additionally, theformation energies imply that substitution of Cl with F enhanceschemical stability due to high electronegativity of fluorine. Figure 1fillustrates the elemental dependence of halogen substitution on theinterfacial energy between c-Li2OHCl0.875X0.125(100) and LLTO(100).As expected, F-substituted c-Li2OHCl exhibits lower interfacial energycompared to other halogen substitutions. These results indicate that Fsubstitution is an effective strategy for achieving both chemical sta-bilization of c-Li2OHCl and optimal interfacial properties in compositesolid electrolytes with LLTO.To address the limitations of pureLi2OHCl, particularly its low ionic conductivity at room temperatureand the cubic-to-orthorhombic phase transition around 345K, weinvestigated the effects of halogen substitution on the structure andionic transport properties of Li2OHCl.MD simulations were conductedin an NVT ensemble with a Nosé-Hoover thermostat. Before conduct-ing the large-scale MD simulations involving over 400 atoms withdurations of 2 ns, both PFP-MD and AIMD (ab initio MD) simulationswere performed to validate their reliability in the Li2OHCl andLi2OHCl0.875F0.125 systems, each consisting of 40 atoms. The corre-sponding initial and final configurations are available in Supplemen-tary Data 3. As shown in Supplementary Fig. 2, the mean-squareddisplacement (MSD) of Li and H during these MD simulations exhibitsresults consistent with those previously reported by Wang, F. et al.23andHoward, J. et al.24. No evidenceof free diffusion ofH⁺wasobservedin this study, confirming that ionic conductivity is solely derived fromLi ions. For this study, PFP-MD was employed for the quantitativeevaluation of interface stability and Li-ion diffusion effects, with itsreliability verified through comparisons with corresponding totalenergies obtained via DFT.Figure 2a shows the dependence of halogen substitutionon the Li-ion diffusivity of c-Li2OHCl0.875X0.125 based on MD simulations, as wellas the volume fraction of available Li⁺ sites determined using the bondvalence sum (BVS) method25,26. The corresponding halogen-substituted structures are available in Supplementary Data 4. Theionic conductivity of c-Li2OHCl0.875F0.125 was calculated to be higherthan that of other halogen substitutions, even though the latticeparameter of c-Li2OHCl0.875F0.125 is smaller compared to those of otherhalogen substitutions (Fig. 1e). Generally, a decrease in lattice para-meter reduces the available volume for ion passage, which typicallyleads to increased activation energy and decreased ion diffusivity27.However, despite a reduced cell volume following F substitution, anincreased volume fraction of available Li⁺ sites was observed (Fig. 1e).Figure 2b illustrates three distributions related to Li-ion movement:volumetric distribution for available Li+ sites from the BVS analysis,energy isosurfaces from the Bond Valence Energy Landscape (BVEL)24,and probability density distributions derived from MD simulations.Geometrical rearrangements around F sites with smaller ionic radiuscreate additional space in the Cl-based framework; conversely, Br or Isubstitutions with larger ionic radii reduce this effect by decreasingavailable volumes for Li+. The distribution of available volumes closelymatches migration pathways identified in both BVEL isosurfaces andMD simulations (Fig. 2b). Detailed results of MD simulations regardingF substitution are provided in Supplementary Fig. 3, and the corre-sponding initial and final configurations are provided in Supplemen-tary Data 5. An increase in F dopant concentration leads to a reductionof the activation energy for Li transport, accompanied by a decrease inthe lattice parameter (Supplementary Fig. 3j). Therefore, F-doped c-Li2OHCl emerges as a promising candidate for an effective lattice-matched system in conjunction with LLTO as a composite solidelectrolyte.We then confirmed the phenomenon of the cubic-to-orthorhombic phase transition in F-substituted Li2OHCl for the roomtemperature operation of ASSBs. The phase transition of Li2OHCl isdriven by changes in OH- orientation26–28. At low temperatures,o-Li2OHCl adopts an ordered structure where Li ions occupy specificWyckoff sites (2a and 2b) within the Pmc21 space group, with localcharge compensation achieved through OH- alignment towardvacancies (Supplementary Fig. 4). In contrast, at higher temperatures,increased thermal vibrations promote site hopping and rotationalmotion ofOH- ions26–28. This results in a disordered c-Li2OHCl structurewhere approximately 66.6% of 3c Wyckoff sites (Pm-3m) are occupiedby Li ions while others remain vacant. This transition involvespendulum-like rotationalmotion of OH- ions coupledwith enhanced Limigration. Figure 2c shows the temperature dependence of MSD of Liions (Li-MSD) calculated from MD simulations with and without Fsubstitution. In pure Li2OHCl, Li is localized at the 2a and 2b sites, withno hopping observed to 3c sites; consequently, Li-MSD remainsunchanged up to approximately 325 K. A discontinuity in Li-MSD isobserved between 325 K and 350K as Li begins occupying the 3c sites.Our simulation results (Supplementary Fig. 3, 4 and SupplementaryData 6) confirm random angular variations in OH- orientations abovethe phase transition temperature, consistentwith previous studies26–28.In the case of F substitution, continuous changes in Li-MSD areobserved even at temperatures as low as 275 K (Fig. 2c). Li occupies allthree sites, 2a, 2b, and 3c, in the cubic arrangement. As shown inSupplementary Fig. 4, the bond length of Li-F is shorter than that of Li-Cl, resulting in the displacement of Li+ from their original orderedpositions in o-Li2OHCl to disordered positions associated with dis-turbances in OH- orientation. Radial distribution function (RDF) of theoptimized Li2OHCl0.875F0.125 structure closely resembles that of thec-Li2OHCl (Supplementary Fig. 4c), indicating that F substitutionmodifies the atomic arrangement toward the disordered cubicstructure.Li-ion transport within a Li2OHCl/LLTO interface model wasinvestigated using MD simulations. The corresponding initial and finalconfigurations are provided in SupplementaryData 7. The LLTOmodelgenerated using REMC exhibits a two-dimensional structure withalternating La-rich and La-poor layers, where the La-rich layers tend tohinder Li-ion conduction16,17. In this model, the La layers are alignedArticle https://doi.org/10.1038/s41467-025-62860-1Nature Communications | (2025) 16:7372 3www.nature.com/naturecommunicationsalong the x-z plane, restricting Li+ transport to two dimensions withinthe x-z plane and effectively prohibiting conduction in the y-direction(Fig. 2d). MD simulations were performed at temperatures below800K to prevent an unsuitable rare event of O-H bond breakage at theLi2OHCl(100)/LLTO(100) interface. Figure 2d shows a reconstructedisosurface of the Li-ion probability density distribution observed dur-ing MD simulations of the Li2OHCl(100)/LLTO(100) interface model,visually illustrating the diffusion trajectory of Li ions. Li-ion transportoccurs actively within both the bulk regions of Li2OHCl and LLTO, aswell as across their interface, as illustrated by the BVEL isosurface inSupplementary Fig. 5. Figure 2e compares the temperature depen-dence of Li-ion diffusivity for the Li2OHCl(100)/LLTO(100) interfacemodel with those observed in c-Li2OHCl and LLTO bulk crystals. Li-iondiffusivity for the interface model lies between those of each bulkmaterial. The ionic conductivity contribution of bulk LLTO is promi-nent in the x-direction where LLTO forms a continuous pathway,whereas its contribution is negligible in the y-direction due to low Li-ion conductivity along this axis. In contrast, in the z-direction, whereLi2OHCl and LLTO are arranged with equal thickness under periodicboundary conditions, the bulk ionic conductivity (σ) contributions canbe expressed as 1/σtotal = 1/2σLi2OHCl + 1/2σLLTO. Consequently, Li2OHCl,which exhibits lower ionic conductivity, plays a significant role indetermining overall ionic transport properties along the z-direction.To validate this observation, z-directional diffusivity was analyzed andfound to be nearly identical to that in bulk-phase c-Li2OHCl crystals(Fig. 2e), suggesting that ion diffusion in bulk-phase c-Li2OHCl is likelythe rate-limiting step for total conductivity in the z-direction. Figure 2fshows the F concentration dependence of interfacial ionic con-ductivity in the z-direction, determined by MD simulations for thec-Li2OHCl1-xFx/LLTO interfaces. The ionic conductivity across theinterface increases with increasing F content, while F substitution alsoreduces the activation energy (Supplementary Fig. 5d). The enhance-ment of ionic conductivity across the interfaceby F substitution canbeattributed to two factors: (1) the intrinsic improvement in the con-ductivity of bulk c-Li2OHCl1-xFx due to extended migration pathways,and (2) improved interfacial lattice matching with LLTO achievedthrough tuned lattice parameters. These findings demonstrate thathalogen substitution, particularly fluorination, effectively addresseskey challenges associated with Li2OHCl and enables its seamless inte-gration into high-performance composite solid electrolytes suitablefor ASSBs.Synthesis and analysis of Li2OHCl1-xFx-LLTO composite solidelectrolytesComposite solid electrolytes of Li2OHCl1-xFx-LLTO were synthesized.The detailed synthesis procedures are described in the Method sec-tion. Briefly, the composite solid electrolytes were prepared by hotpressing Li2OHCl1-xFx with LLTO powders. Infiltration of LLTO withmolten Li2OHCl1-xFx achieved a pellet density of 99%. Adding LLTO as afiller29 reduces the fluidity of the molten Li2OHCl0.9F0.1, minimizingleakage and facilitating pellet formation through hot pressing. Thephotographs of Li2OHCl0.9F0.1 and Li2OHCl0.9F0.1-LLTO pellets kept atbba cc-Li2OHClLLTODiffusion of Li+zyx-15-14-13-12-11-10ln D (cm2 s-1)-15-14-13-121.2 1.25 1.3 1.35 1.4 1.45 1.5 1.55 1.6 1.65ln D (cm2 s-1)1000/T (K-1)LLTOc-Li2OHCl/LLTOc-Li2OHClc-Li2OHCl/LLTO,z-directionc-Li2OHCl,z-directioned f00.20.40.60.811.21.41.61.8250 300 350 400Li-MSD (Å2 )Temperature (K)00.511.522.533.544.50 0.0625 0.125 0.1875 0.25Li⁺diffusivity along z-direction (10-6cm2s-1)x in c-Li2OHCl1-xFxOH orderingPendulum-like OH rotationLi2OHClLi2OHCl0.875F0.125OH12.51313.51414.51515.51616.500.20.40.60.811.21.41.61.82F Cl Br ILi diffusivity (10-6cm2s-1)Volume fraction for available Li+ sites (%)Available Volume for Li+Li+isosurface (BVEL)Li+isosurface (MD)La layer40020vol%HClTiLaOLiFig. 2 | Li-ion dynamics of lattice-matched c-Li2OHCl1-xFx-LLTO composite solidelectrolytes based onMD simulations. a Dependence of halogen substitution onthe Li-ion diffusivity of c-Li2OHCl0.875X0.125 (X = F, Cl, Br, and I) calculated from Li-MSD, as well as the volume fraction of available Li+ sites determined using the BVSmethod. b Visualized distributions related to Li-ion movement: volumetric dis-tributionof available Li+ sites derived fromBVS shown inblue; energy isosurface forLi+ obtained from the BVEL shown in purple; and probability density distribution ofLi+ obtained from MD simulations shown in yellow. Isosurfaces from MD simula-tions are plotted at an isovalue of 0.005 au. c Temperature dependence of Li-MSDwith and without F substitution. d Reconstructed isosurface illustrating the prob-ability density distribution of Li ions duringMDsimulations. Isosurfaces are plottedat an isovalue of 0.0001 au. The isosurfaces representing Li ions are shown inyellow. e Temperature dependence of Li-ion diffusivities in LLTO, c-Li2OHCl, andLi2OHCl(100)/LLTO(100), as well as along the z-direction in c-Li2OHCl, andc-Li2OHCl(100)/LLTO(100). f F concentration dependence of interfacial ionic con-ductivity in the z-direction for the c-Li2OHCl1-xFx/LLTO interface. Source data areprovided as a Source Data file.Article https://doi.org/10.1038/s41467-025-62860-1Nature Communications | (2025) 16:7372 4www.nature.com/naturecommunicationsdifferent temperatures are shown in Fig. 3a. In the case of the LLTOcomposite, although Li2OHCl0.9F0.1 melts, the pellet structure remainsintact. The experimental results demonstrate that Li2OHCl1-xFx exhibitshigh wettability with LLTO, as shown in Fig. 1c, f, and indicate thatLi2OHCl1-xFx/LLTO has low interfacial energy. However, because thisprocess utilizes the Li2OHCl1-xFxmelt, even slight temperature changessignificantly impact the internal structure of the composite pellet(Fig. 3b). By processing at 548 K (corresponding to the melting onsettemperature), a dense and well-infiltrated composite with no distinctgrain boundaries in the Li2OHCl0.9F0.1 region was obtained. At 543 K(below its melting onset temperature), both Li2OHCl0.9F0.1 and LLTOparticles were observed. At 553 K (above the melting point), excessivemelt was generated, leading to leakage caused by a decrease in visc-osity, which resulted in the formation of cavities within the compositematerial.Figures 3c, d show the XRD results and enlarged patterns forLi2OHCl, Li2OHCl-LLTO, Li2OHCl0.98F0.02-LLTO, and Li2OHCl0.9F0.1-LLTO composite solid electrolyte pellets. Li2OHCl exhibits an orthor-hombic phase at room temperature as expected. In the Li2OHCl-LLTOcomposite solid electrolyte, both orthorhombic and cubic phasescoexist, unlike standalone Li2OHCl. With increasing F dopant con-centration, the orthorhombic phase decreases while the cubic phaseincreases in Li2OHCl1-xFx. The cubic/orthorhombic ratios for x =0,0.02, 0.05, and 0.1 were calculated to be 59, 91, 95, and 100%,respectively, using Rietveld analysis (Supplementary Fig. 6). Thec-Li2OHCl1-xFx peak position, referenced to the LLTO(110) peak at32.8°, shifts toward higher angles for x =0.1 compared to x = 0.02,indicating contraction of the crystal lattice. As shown in Fig. 3e andSupplementary Fig. 6c, the lattice parameter decreases with increasingx, particularly at x =0.1, where it contracts to 3.9025 Å, approximately0.2% smaller than that of typical c-Li2OHCl (3.911 Å). This lattice con-traction reduces the lattice mismatch with LLTO to approximately0.8%. The ionic conductivity of the synthesized Li2OHCl0.9F0.1 at 298 Kalso improves to 3.3 × 10-6 S cm-1 from 1.1 × 10-7 S cm-1 in Li2OHCl,demonstrating an inverse correlation with lattice parameters (Fig. 3e).These results align with theoretical calculations in Figs. 1e and 2a,supporting the validity of F substitution effects. It is worth noting thatan increasing trend of small LiF precipitates is observed at higher Fcubicorthorhombic3.93.9013.9023.9033.9043.9053.9063.9070 0.02 0.04 0.06 0.08 0.1Li+conductivity (S cm-1)Lattice Parameter ( Å)x in Li2OHCl1-xFx10-510-610-7Li2OHCl0.9F0.1-LATPLi2OHCl0.9F0.1-LLZOLi2OHCl0.9F0.1Li2OHCl0.9F0.1-LLTO02460 2 4 6-Z'' (kΩ)Z' (kΩ)298 K31.5 32 32.5 33 33.52θ (degree)O+*O+O*+* ** * *aaLi2OHCl0.9F0.1pelletLi2OHCl0.9F0.1-LLTOpellet298 K298 K 548 K548 Kec d*20 25 30 35 40 45 502θ (degree)*Li2OHCl pelletLi2OHCl0.98F0.02-LLTO pellet298 KLi2OHCl0.9F0.1-LLTO pelletLi2OHCl-LLTO pellet+O OOOOOOOO++ ++O OOOOOOOO++ +* * *+O OOOOOOOO++ +** ******* * * ***** **o-Li2OHCl (Pmc21)c-Li2OHCl (Pm-3m)LLTO (P4/mmm)+O*Intensity (arb. units)fb 543 K 548 K 553 Kg012340 1 2 3 4-Z'' (kΩ)Z' (kΩ)measuredfittedR0 R1 R2CPE2CPEwCPE1Fig. 3 | Characterization of Li2OHCl1-xFx-LLTO composite solid electrolytes.a, Photographs of Li2OHCl0.9F0.1 and Li2OHCl0.9F0.1-LLTOpellets stored at differenttemperatures. b Cross-sectional SEM images of Li2OHCl0.9F0.1-LLTO pellets syn-thesized at 543, 548, and 553 K. The scale bars in b are 10μm. c XRD patterns atroom temperature for the synthesized Li2OHCl, Li2OHCl-LLTO, Li2OHCl0.98F0.02-LLTO, and Li2OHCl0.9F0.1-LLTO composite solid electrolytes. d Enlarged view ofc with Rietveld analyses included for detailed structural assignments.eRelationshipbetweenchanges in latticeparametersdue to Fdoping in Li2OHCl1-xFand the measured Li ionic conductivity of Li2OHCl1-xFx at 298K. f Nyquist plots ofimpedance at 298K for Li2OHCl0.9F0.1, Li2OHCl0.9F0.1-LLTO, Li2OHCl0.9F0.1-LLZO,and Li2OHCl0.9F0.1-LATP composite solid electrolyte pellets. g Nyquist plot, fittedcurve, and the corresponding equivalent circuit of Li2OHCl0.9F0.1-LLTO compositesolid electrolyte pellets. Source data are provided as a Source Data file.Article https://doi.org/10.1038/s41467-025-62860-1Nature Communications | (2025) 16:7372 5www.nature.com/naturecommunicationsdopant concentrations, as shown in Supplementary Fig. 7. Theseobservations suggest that a composition around x = 0.1 likely repre-sents a practical solubility limit.Figure 3f shows Nyquist plots of impedance at room temperaturefor c-Li2OHCl0.9F0.1, Li2OHCl0.9F0.1-LLTO, Li2OHCl0.9F0.1-LLZO, andLi2OHCl0.9F0.1-LATP composite solid electrolyte pellets. TheLi2OHCl0.9F0.1-LATP composite solid electrolyte exhibited higherresistance than c-Li2OHCl0.9F0.1, indicating that the high ionic con-ductivity of LATPwasnot effectively realized. Thismaybe attributed tothe evidently high interfacial energy, as shown in Fig. 1c, which sug-gests that an appropriate interface for Li-ion transport could not beformed. In contrast, it was found that composites with LLTO or LLZOresulted in lower resistance compared to c-Li2OHCl0.9F0.1, with a sig-nificant reduction in resistivity observed for the c-Li2OHCl0.9F0.1-LLTOcomposite solid electrolyte. Notably, two distinct semicircles wereclearly observed for both composites with LLTO and LLZO in Fig. 3f. Asshown in Fig. 3g and Supplementary Fig. 8, equivalent circuit modelanalysis suggests that the first and second semicircles represent theionic conductivity at the Li2OHCl0.9F0.1-LLTO interface (1.7 × 10-4 S cm-1)and the bulk ionic conductivity of Li2OHCl0.9F0.1 (9.7 × 10-6 S cm-1),respectively. The low resistance of the Li2OHCl0.9F0.1/LLTO interface,attributed to its energetically stable lattice-matched structure, wascomparable to the resistance observed at sintered LLTO grainboundaries27. The total ionic conductivity of the c-Li2OHCl0.9F0.1-LLTOcomposite solid electrolyte was calculated to be 5.0 × 10-5 S cm-1 (Sup-plementary Fig. 8). This value is approximately one order ofmagnitudehigher than that of c-Li2OHCl0.9F0.1, facilitating ASSB operation atroom temperature.Pressure-assisted melt-infiltrated ASSBs with conventional LIBelectrodesASSBs were fabricated using the Li2OHCl0.9F0.1-LLTO composite solidelectrolyte via melt infiltration into conventional LIB electrodes. Atape-cast NCA (LiNi0.8Co0.15Al0.05O2) secondary particle positive elec-trode and a tape-cast Li4Ti5O12 (LTO) secondary particle negativeelectrode were utilized, with relatively higher calendar pressingapplied to induce significant microcracks in the active materials,thereby accelerating melt infiltration. A compacted Li2OHCl0.9F0.1-LLTO composite solid electrolyte sheet was sandwiched between thepositive and negative electrodes, and ASSBs were fabricated by heat-ing and melt-infiltrating the solid electrolytes. Xiao et al. previouslyreported a thermal melt infiltration method11. As discussed above,handling Li2OHCl melt poses significant challenges due to leakageduring processing. In the fabrication of ASSBs, issues arise fromincomplete infiltration into fine voids within the electrode, as well asthe potential for short circuits between the positive and negativeelectrodes. In contrast, with Li2OHCl0.9F0.1-LLTO composite solidelectrolytes, stable pressure-assistedmelt infiltration30 by hot pressingis feasible. As shown in Fig. 4a, LLTO functions as a separator filler toprevent short circuits when pressurization. Figure 4b shows SEMimages of the cross-section of the negative electrode after thermal andpressure-assisted melt infiltrations. In thermal melt infiltration,Li2OHCl0.9F0.1 infiltrates into the negative electrode; however, manycavities smaller than 200 nm are observed. In contrast, pressure-assistedmelt infiltration achieves higher penetration, effectively fillingboth the interior of the electrode and the secondary particles ofnegative electrode active materials. External pressure allows moltenThermal melt infiltration Pressure-assisted melt infiltrationbbcSECross-sectional SEMLi+ Ni+ C+ Al+Ti+ Cl- O- F-ToF-SIMS positive mapping ToF-SIMS negative mappingAleSEM O CTi Cl LaSEAldaSE separatorPositive electrodeNegative electrodeLLTOInfiltrationInfiltrationHot pressingHot pressingCross-sectional-STEM EDX mappingNi Cl CPositive electrodeNegative electrodeFig. 4 | Melt infiltration of Li2OHCl0.9F0.1-LLTO composite solid electrolyte intoconventional tape-cast LIB electrodes. a Schematic illustration of the cross-sectional structure of ASSB with Li2OHCl0.9F0.1-LLTO and conventional tape-castLIB electrodes. b Cross-sectional SEM images of LTO negative electrodes withLi2OHCl0.9F0.1-LLTO composite solid electrolyte prepared by thermal and pressure-assistedmelt infiltrations. c SEMandToF-SIMSmapping imagesof theNCApositiveelectrode cross-section after pressure-assistedmelt infiltration. d STEM and STEM-EDX mapping image of an NCA positive electrode after pressure-assisted meltinfiltration; Green, blue, and red represent Ni, Cl, and C, respectively. e Three-dimensional SEM cross-sectional image and the EDX mappings of Li2OHCl0.9F0.1-LLTO composite solid electrolyte prepared by pressure-assisted melt infiltration.The scale bars in b, c, d, and e are 2, 10, 2, and 10μm, respectively.Article https://doi.org/10.1038/s41467-025-62860-1Nature Communications | (2025) 16:7372 6www.nature.com/naturecommunicationssolid electrolytes to penetrate delicate and complex pathways. Thepacking density of the entire ASSB after pressure-assisted melt infil-tration reached approximately 96%, compared to 91% for thermalmeltinfiltration. Our investigation revealed that pressure-assisted meltinfiltration provides improved penetration compared to conventionalthermal methods and allows for effective infiltration into complexelectrode geometries. This increased packing density is crucial formaximizing ionic conductivity throughout the composite electrodesystem. However, as shown in Fig. 3b, this process has a very narrowtemperature range: above 548K, Li2OHCl0.9F0.1 begins to melt, andunder pressure, molten salt may flow outward from the electrode.Accurate temperature control is essential for stable penetration intothe electrode. At an optimized temperature for pressure-assisted meltinfiltration (548K) and pressing pressure (5MPa), controlled meltingoccurs on the Li2OHCl0.9F0.1 surface whileminimizing external leakageof molten salt and promoting efficient penetration into electrodevoids. Notably, these optimal values may vary depending on equip-ment or material compositions.Figure 4c shows an SEM image and time-of-flight secondary ionmass spectrometry (ToF-SIMS) mapping images of the positive elec-trode cross-section after pressure-assisted melt infiltration. It wasconfirmed that Li2OHCl0.9F0.1, serving as the ionic pathway, penetrateddeeply into the densely compacted electrode. Figure 4d shows scan-ning transmission electron microscopy (STEM) and energy-dispersiveX-ray spectroscopy (STEM-EDX) images of the positive electrodeactive material region, where Li2OHCl0.9F0.1 effectively filled micro-cracks in the active material, providing a path for Li-ion conduction.This infiltration structure was found to replicate the behavior of liquidelectrolytes by creating efficient pathways for Li ions. Three-dimensional (3D) reconstruction images (Fig. 4e), generated fromSEM cross-sectional and SEM-EDX mapping obtained through con-tinuous focused ion beam (FIB) processing, suggest that the interior ofthe tape-cast negative electrode is densely filled with solid electrolyte,exhibiting an electrodeporosity of less than 1%. This lowporosity helpsmaintain the mechanical integrity of the electrode while also enhan-cing ionic conductivity within the electrode. The curvature of theLi2OHCl0.9F0.1 ion pathway, calculated from these 3D images, is 2.29,comparable to an equivalent value of 2.11 for ionic pathways in liquidLIB electrodes as evaluated by ionic conductivity. Additionally, smallerLLTO particles were observed to be dispersed throughout the elec-trode interior alongside infiltrated Li2OHCl0.9F0.1 (as shown in the Lamapping in Fig. 4e). This phenomenon is likely attributed to capillaryeffects, indicating transport of small LLTO particles during the infil-tration of Li2OHCl0.9F0.1 into complex pathways. Such dispersion ofLLTO is presumed to enhance ionic conductivity within the electrodeinterior.The electrochemical properties of the pressure-assisted melt-infiltrated ASSB coin cells were evaluated. Figure 5a, b, and c show aphotograph of the ASSB coin cell, a detailed schematic of its internalstructure, and a processed image of the ASSB junction obtainedthrough SEM-EDX imaging, respectively. The highly infiltrated solidelectrolyte in the electrode enables the ASSB junction to be tightlysolidified, eliminating the need for excessive cell pressure. Forinstance, the ASSB junction can operate by simply embedding it in anexternal casing (Fig. 5a, b) and has been verified to retain functionalityeven when cut into smaller pieces. As shown in Fig. 5c, the infiltratedsolid electrolyte allows for the use of tape-cast electrodes designed forconventional LIBs, significantly reducing electrode costs. This design isstraightforward and accommodates various cell configurations. Com-mon winding or stacking structures used in liquid LIBs can also beimplemented.Figure 5d shows the Nyquist plots of the ASSBs before charging at298K utilizing composite solid electrolytes containing LLTO, LLZO,and LATP. In the ASSB configuration, tape-cast LIB electrodes werecombined with relatively higher calendar pressing to inducemicrocracks and enhance specific surface area. Consequently, the cellresistance was reduced to approximately one-tenth that of the com-posite solid electrolyte with Pt-sputtered electrodes shown in Fig. 3f.Among these composites, LLTO demonstrated a relative low cellresistance, while LATP exhibited a relatively high cell resistance, con-sistent with the trends observed in Fig. 3f.Figure 5e shows the initial charge-discharge characteristics of theASSBs at 298K using composite solid electrolytes containing LLTO,LLZO, and LATP. The low resistivity of the Li2OHCl0.9F0.1-LLTO com-posite solid electrolyte significantly enabled the room-temperatureoperation of this ASSB. In comparison to LLTO, incorporating otheroxide solid electrolytes as filler materials increased cell resistivity(Fig. 5d), leading to higher overpotentials in the ASSBs and a reductionin achievable charge-discharge capacity (Fig. 5e).Figure 5f shows the cycle performance of the ASSB usingLi2OHCl0.9F0.1-LLTO composite solid electrolyte at 333 K. The capacityretention of 92% was maintained after 100 cycles. A slight increase inoverpotential (0.1 V for both charge and discharge) was observed inthe charge-discharge curves after 100 cycles (Supplementary Fig. 9a).The capacity retention was recovered to 99.4% through low-rateoperation (0.067 C), indicating limited degradation of the activematerials during the cycling test. The Nyquist plots before and aftercycling are shown in Supplementary Fig. 9c. Based on the equivalentcircuit model analysis of the impedance components (SupplementaryFig. 9d and 9e), the ohmic resistance component, which is attributedto electronic conduction within the electrode and at the electrodeinterface, remains stable. However, both the high-frequency and low-frequency components, which are likely associated with chargetransfer resistance at the negative electrode and positive electrodeinterfaces, are observed to increase. Notably, the second semicircleexhibits a significant increase. The impedance value corresponding tothe peak of this second semicircle at 1 kHz demonstrates a slight linearupward trend with repeated charge-discharge cycles (SupplementaryFig. 10b). Supplementary Fig. 10 shows the dependence of the capacityretention and interface resistivity (impedance value at 1 kHz) of theASSB cell on calendar pressing of the positive and negative electrodes.The effect of calendar pressing on the porous negative electrode wasnegligible, whereas its impact on the interface resistivity at the positiveelectrode was significant for both capacity retention and interfaceresistivity. Supplementary Fig. 10c shows theNyquist plots of the ASSBat the 100th cycle with and without calendar pressing of the positiveelectrode. An increase in the ohmic resistance component wasobserved, suggesting delamination at the positive electrode interface.Furthermore, an increase in the second semicircle was confirmed,indicating an increase in charge transfer resistance at the positiveelectrode interface. The high-pressure pressing induces cracks indenseNCA secondary particles, promoting solid electrolyte infiltrationinto these cracks (Fig. 4d). For liquid electrolytes, increased reactionsurface area due to cracks within active material can cause surfacedegradation, leading to reduced cycle life and safety performance31.However, for the infiltrated solid electrolyte, it is speculated thatcracks withinNCA secondary particles contribute to the formation of arobust self-binding electrode via an anchoring effect while also redu-cing resistance through enhanced solid electrolyte penetration. Thisphenomenon likely stabilizes interface resistance and improves cycleperformance.To reduce resistivity and enhance the adaptability of ASSBs, theuse of smaller NCA positive electrode particles with a higher surfacearea was demonstrated. The rate performance at 298K for an ASSBemploying smaller NCA particles (D50: 5 µm) and a thinner solidelectrolyte layer (thickness: 50 µm) is shown in Figs. 5g and 5h. Stableoperation was achieved at rates up to 0.5 C, whereas ASSBs with largerNCA particles (D50: 15 µm) exhibited a maximum rate of 0.2 C (Sup-plementary Fig. 11). These results suggest potential for furtherimproving electrochemical performance by optimizing electrodeArticle https://doi.org/10.1038/s41467-025-62860-1Nature Communications | (2025) 16:7372 7www.nature.com/naturecommunicationsconfigurations, as well as enhancing the ionic conductivity of Li2OHCland reducing the interfacial resistivity at the Li2OHCl/LLTO interfacethrough increased fluorine doping.In summary, while composite solid electrolytes aim to effectivelyleverage the advantages of their constituent materials, our studydemonstrates that lattice matching and interfacial stability at theirinterfaces have a profound impact on ionic conductivity and electro-chemical performance. Moreover, well-designed interface modelingbased on structurally compatible materials can enhance latticematching and interfacial stability, thereby enabling more stableoperation of ASSBs. Additionally, we havedemonstrated that pressure-assisted melt infiltration using molten salts plays a pivotal role inachieving such optimized interfacial structures. This underscores theimportance of exploring molten salt-based solid electrolytes as a keystrategy for realizing all-solid-state batteries that outperform conven-tional liquid LIBs.Building upon these findings, the combination of lattice-matchedcomposite solid electrolytes and pressure-assisted melt infiltration inASSBs demonstrates several attributes and performance enhance-ments, including stable structural integrity, improved cycling stability,flexible design options, and room-temperature operation. Theseadvancements represent a contribution to progress in the develop-ment of advanced energy storage systems, providing useful insightsfor future researchwhile supporting furtherwork in high-performanceall-solid-state batteries.MethodsCalculation of reciprocal lattice projectionsRecipro32, a crystallographic software specialized in simulating dif-fraction patterns and reciprocal lattice projections, was used to cal-culate reciprocal lattice projections for several solid electrolytecrystals. The structures considered were Pm-3m for Li2OHCl26,P4mmm for Li0.31La0.56TiO333, Ia-3d for Li7La3Zr2O1234, and R-3c forLi1.3Al0.3Ti1.7(PO4)335. The reciprocal lattice projections were generatedusing the Recipro diffraction simulator by overlaying individual reci-procal lattice patterns of Li2OHCl and the other solid electrolytes. The00.511.522.530 50 100 150 200Voltage (V)Specific capacity (mAh g-1)298 Kw/ LLTOw/ LLZOLi2OHCl0.9F0.1w/ LATP10 mA g-1 (0.067C)aa bcdgfehTape-cast positive electrode with infiltrated SEactive material: NCA secondary particlesbinder: PVdFconductive additive: VGCFLi2OHCl0.9F0.1/LLTO lattice-matched SETape-cast negative electrode with infiltrated SEactive material: LTO secondary particlesbinder: PVdFconductive additive: VGCF11.21.41.61.822.22.40 100 200Voltage (V)Specific Capacity (mAh g-1)298 K0.1C0.2C0.3C0.5C1C8085909510010511011512050607080901001100 20 40 60 80 100 120Coulomb Efficiency (%)Capacity Retention (%)Cycle number333 K50 mA g-1 (0.3C)10 mA g-1 (0.067C)95969798991001011021031041050204060801001201401601800 5 10 15 20Coulomb Efficiency (%)Specific Capacity (mAhg-1)Cycle number0.1C 0.2C0.3C0.5C1C0.1C298 K020040060080010000 200 400 600 800 1000-Z’' (Ω)Z' (Ω)w/ LLTOw/ LLZOLi2OHCl0.9F0.1w/ LATP298 KFig. 5 | Electrochemical performance of LTO | Li2OHCl0.9F0.1-LLTO | NCA ASSBcoin cells fabricated by pressure-assisted melt infiltration. a Photograph of anASSB coin cell. b Detailed schematic of the internal structure of the ASSB coin cell.c Processed image of the ASSB junction obtained through SEM-EDX imaging. d, eNyquist plots of impedance and initial charge-discharge characteristics at 298K forASSBs with different oxide solid electrolyte fillers. f Cycle performance of the ASSBusing Li2OHCl0.9F0.1-LLTO composite solid electrolyte at 333K under a charging/discharging current of 50mAg-1 (corresponding to 0.3 C). The last data point after100 cycles was obtained using a charging/discharging current of 10mAg-1 (corre-sponding to0.1 C) tominimize the effect of cell resistance.g,hRate performanceat298K for ASSBs utilizing smallNCAparticles (D50: 5 µm)and a thin solid electrolytelayer (thickness: 50 µm), selected to reduce ionic resistance and enhance rateperformance. Source data are provided as a Source Data file.Article https://doi.org/10.1038/s41467-025-62860-1Nature Communications | (2025) 16:7372 8www.nature.com/naturecommunicationslattice plane indices of each structure were identified based on ener-getically favorable surfaces calculated via DFT, as referenced instudies36–38.Preparation of Li2OHCl1-xFx, LLTO, LLZO, LATP, and thoseinterface models with Li2OHCl1-xFxThe stable configurations of randomly structured crystals ofLi2OHCl1-xFx (240 atoms), LLTO (234 atoms), LLZO (372 atoms), andLATP (440 atoms) (shown in Supplementary Fig. 2a and Supplemen-tary Data 1) were obtained through replica exchange Monte Carlo(REMC)19 calculations using the PFP universal neural-network potential(version 5.0.0) provided by Matlantis20, which was integrated into theAtomic Simulation Environment (ASE)39. This machine learning surro-gate model for ab initio calculations supports the elements utilized inthis study, which include H, Li, O, F, Cl, Ti, and La. The reliability of thePFP during these simulations was verified by comparing the corre-sponding energy calculations and RDFs obtained via DFT, with a MeanAbsolute Error (MAE) of less than 10-3eV/atom, as shown in Supple-mentary Fig. 2. The Metropolis sampling was performed on systemreplicas at different temperatures in parallel. The selected tempera-tures ranged from 300 to 2100K to enable enhanced sampling andexploration of diverse energy landscapes.The interface models were prepared by combining each struc-ture and enlarging the supercell size to ensure compatibility.Li2OHCl1-xFx/LLTO (474 atoms), Li2OHCl1-xFx/LLZO (5348 atoms), andLi2OHCl1-xFx/LATP (5732 atoms) were successfully constructed(Supplementary Data 2).DFT calculations were performed using the Vienna ab initioSimulation Package (VASP)40,41. The Perdew-Burke-Ernzerhof (PBE)exchange-correlation functional42 was combined with the projectoraugmented wave (PAW) method43,44, a plane-wave basis set with anenergy cutoff to 520 eV. Long-range dispersion forces were accountedfor using the Grimme DFT-D3 method45. Due to the large supercellemployed in these simulations, only Γ-point sampling was performed.The following valence electron configurations were adopted: 1s1 for H;1s22s1 for Li, 2s22p4 for O; 3s23p5 for Cl; 3d34s1 for Ti; and5s25p65d16s2 for La.The VESTA software package46 was used to visualize crystalstructures and isosurfaces.MD calculations for Li2OHCl1-xFx, LLTO and Li2OHCl1-xFx-LLTOinterface modelsMD simulations were performed using the PFP provided by Matlantis.The reliability of the PFP during these simulations was verified bycomparing the corresponding energies, forces, RDFs, and MSDs withthose obtained via DFT, as shown in Supplementary Fig. 2. To furtherensure the reliability of PFP-MD, AIMD simulations were performed onLi2OHCl and Li2OHCl0.875F0.125 systems, each consisting of 40 atomsover a duration of 20ps (40,000 steps). All simulations were con-ducted with a timestep of 0.5 fs to account for the movement ofhydrogen atoms in the simulation. The temperature was controlledusing the Nosè-Hoover thermostat for maintaining equilibrium atconstant temperature and Parrinello-Rahman barostat for enablingisotropic pressure control in phase transition analysis. To analyze thephase transition behavior of Li2OHCl1-xFx, MD simulations were per-formed using both the NVT and the NPT ensembles. Analyses of Li andOH ion dynamics were conducted in the NVT ensemble for 2 ns acrossa temperature range of 250–400K, with increments of 25 K, to capturedetailed dynamic behavior near and beyond the phase transitiontemperature (345 K). Lattice parameter analyses during the phasetransition were conducted in the NPT ensemble, with the temperatureincrementally increased by 5 K at each step over a duration of 20ps.The calculated lattice parameters at each step were obtained byaveragingmeasurements taken after excluding the initial 10 ps of eachstep. The simulation temperature was determined from the time-averaged kinetic energy. Additionally, a combination of the Berendsenthermostat and Berendsen barostat was also employed to verifyreproducibility of phase transition behavior under a temperaturegradient.MD simulations to study Li-ion diffusion were performed usingthe NVT ensemble with varying temperatures controlled by the Nosé-Hoover thermostat using the PFP. O-H bond breakage was observed inthe Li2OHCl/LLTO interface at temperatures above 825 K; this phe-nomenon could affect structural stability and ionic conductivity pre-dictions under such high temperature conditions; therefore,temperatures of 650, 675, 700, 725, 750, and 775 K were selected toevaluate ion dynamics while maintaining the structural integrity of thesimulated systems. The simulationswere run for 2 ns. The isosurface ofLi-ion probability density distribution was visualized using atomictrajectories generated from the MD simulations. The self-diffusioncoefficients of Li ions were calculated from theMSD using the Einsteinrelation. The activation energies for Li-ion diffusion were calculatedusing Arrhenius plots of the self-diffusion coefficients. The corre-sponding initial and final snapshots of MD trajectories, saved as CIFfiles, are provided in the Supplementary Data.BVEL calculations for Li ion migration pathwayBond Valence Energy Landscape (BVEL) calculations were performedusing the BondStr24 program with the softBV23 model, which providesan enhanced description of ionic interactions compared to traditionalBVS models, to simulate Li ion migration pathways and evaluatemigration barrier heights. Structural models optimized via PFP andDFT were used as input structures. BVELs of Li ions were calculated ona 3D point grid with a resolution of 0.1 Å. The ion migration pathwaywas determined by calculating an isosurface using graph theory andthe marching cubes algorithm24, which constructs the isosurface as atriangular mesh in the program.Synthesis of composited solid-state electrolytesAll materials used in this study were handled in a dry-room environ-ment with a dew point below 223 K. The synthetic precursors ofLi2OHCl1-xFx were LiOH (anhydrous, 99.9%, Sigma-Aldrich), LiCl(anhydrous, 99.95%, Sigma-Aldrich), and LiF (≥99.99%, Sigma-Aldrich).All precursors were dried at 373 K under vacuum for 24 h prior to theexperiments.LiOH, LiCl, and LiF were mixed at an appropriate molar ratio andground using a ball mill (P-7, Fritsch) at room temperature for 10 h at300 rpm without intermittent rest periods. Throughout the entire ballmilling process, a 45mL zirconia jar and 50g of zirconia balls (3mmdiameter) were used, with approximately 20mL of total powdervolume, corresponding to a ball-to-powder weight ratio of 2:1. Themixture was then heated to 603K for 30min at a ramping rate of 10K/min in a box furnace until it melted into a transparent molten state,ensuring homogeneity of the solid electrolyte composition uponcooling. The molten samples were subsequently cooled to roomtemperature. The resulting white brittle salts were ground into finepowders using amortar and ball milling for 30min at 300 rpmwithoutintermittent rest periods. These steps were repeated three times toobtain Li2OHCl1-xFx solid electrolyte powders.Perovskite Li0.29La0.57TiO316,17 powder (TP-02F, Toho Titanium)and Li0.33La0.55TiO3 powder (Li0.33La0.55TiO3, Toshima Manufacturing)were used as the LLTO materials. The obtained solid electrolyte pow-ders of Li2OHCl1-xFx were mixed with LLTO powder by using ball mil-ling for 30min at 100 rpm without intermittent rest periods. Weightratios of 3:1 (Li2OHCl1-xFx: LLTO) were used for ASSB configurations,and 1:1 for analyses of the composite solid electrolyte itself. Theseratios were optimized to achieve consistent ionic conductivity whilemaintaining structural stability in the composite electrolytes. Addi-tionally, they ensured compatibility and reliable performance in melt-infiltrated electrolytes within electrodes for ASSBs. The mixture ofArticle https://doi.org/10.1038/s41467-025-62860-1Nature Communications | (2025) 16:7372 9www.nature.com/naturecommunicationssolid electrolyte powderswasheated to 558 K for 10min in a preheatedbox furnace until it melted into a whitish molten state. The moltensamples were cooled to room temperature, and the resulting whitebrittle salts were ground into fine powders using a mortar and ballmilling for 10min at 200 rpm without intermittent rest periods. Theseprocesses were repeated three times to obtain Li2OHCl1-xFx-LLTOcomposite solid electrolyte powders.Preparation of active materials and electrodesPositive electrode activematerials: Theprecursors for synthesizingNCA(LiNi0.8Co0.15Al0.05O2) secondary particles were Li2CO3 (99.99%,Kojundo Chemical), NiSO4·6H2O (99.9%, Kojundo Chemical), NaOH(99%, Kojundo Chemical), NH4OH (28.0–30.0%, Sigma-Aldrich) andAl2(SO4)3·18H2O (Sigma-Aldrich). The Ni0.8Co0.15Al0.05(OH)2 precursorwas prepared in a continuously stirred tank reactor via the hydroxideco-precipitation method47. NaOH(aq) solution was used as the basesource for the reaction, while NH3(aq) solution facilitated spherical anddense particle growth. The reactor temperature was maintained at323K, and the contents were stirred at 1200 rpm. A pH controllerautomatically regulated NaOH injection using a peristaltic pump con-nected to the reactor. After 20 h, the hydroxide precursor was rinsedseveral times with water and dried at 393K for 10 h. TheNi0.8Co0.15Al0.05(OH)2 precursor was then mixed with Li2CO3 by ballmilling, with a Li-to-transition metal molar ratio of 1.02 maintainedduring preparation. The mixed powder was heated in a box furnaceunder an oxygen flow at 993K for 20 h at a ramping rate of 10K/min tosynthesize NCA powders. The final NCA powders were mechanicallygroundandpassed through a 53μmsieveprior to electrode fabrication.Negative electrode active materials: The materials used for syn-thesizing Li4Ti5O12 (LTO) secondary particles were Li4Ti5O12 (<200 nm,Sigma-Aldrich) and sucrose (C12H22O11, Kanto Chemical). Sucrose wasdissolved in deionized water to achieve a 0.2wt% solution. LTO pow-der was then suspended in the solution, and the suspension wasultrasonicated for 15min to form a homogeneous dispersion. Thissuspension was processed in a spray dryer to obtain spherical sec-ondary particles48. Sucrose served as both a binder and a carbonsource for the LTO secondary particles. To obtain carbon-coated LTOsecondary particles, calcination under argon flow was performed at973 K for 5 h at a ramping rate of 10K/min. The resulting LTO powderswere mechanically ground and passed through a 53μm sieve beforeelectrode fabrication. Electrode fabrication: NCA powders, LTO pow-ders, vapor-grown carbon nanofibers (VGCF, Showa Denko), andpolyvinylidene fluoride (PVDF; Kureha,molecular weight: >1,000,000)were used to prepare electrodes. The weight ratio of active materials,VGCF and PVDF, was set at 92:5:3 in N-methyl-2-pyrrolidone (NMP) toform a slurry. Single-side-coated positive and negative electrodes wereprepared by casting the slurry onto aluminum foils (thickness: 12μm,purity: >99.5%), followed by drying in an oven at 363 K for 30min. Thecasting was performed using the doctor blade method. Both thepositive and negative electrodes were punched out to a 16mm dia-meter. The electrodes were then calendared under a pressure of20MPa and dried further in a vacuumoven at 393 K for 12 h before cellassembly.Pressure-assisted melt infiltration processAll procedures were conducted in a dry-room environment with a dewpoint below 223 K. Positive and negative electrode discs with a dia-meter of 15mm were prepared. Li2OHCl1-xFx-LLTO composite solidelectrolyte powders were calendared under a pressure of 20MPa toform pellets with a thickness of 120μm. The composite solid electro-lyte pellet was placed between the positive and negative electrodesand heated to 533–553K for 20min while applying hot pressing with apressure of 0-10MPa (P5058, NPa SYSTEM). Careful control of tem-perature and pressure was crucial to prevent melting and leakage ofthe molten solid electrolyte, thereby ensuring cell quality.Fabrication of ASSB coin cellsThe fabricated LTO | Li2OHCl0.9F0.1-LLTO | NCA assembled sampleswere sealed in CR2016 coin cells for electrochemical measurements.The thickness of the composite electrolyte was 100μm, with massloadings of the NCA (D50: 15 µm) positive electrode and LTO negativeelectrode at 10mg cm-2 and 8mg cm-2, respectively. For the rate cap-ability tests, the composite electrolyte thickness was reduced to50 µm, and the mass loadings for the NCA (D50: 5 µm) positive elec-trode and LTO negative electrode were adjusted to 10mg cm-2 and8mg cm-2, respectively.CharacterizationsFE-SEM (S-4800, Hitachi) and FIB-SEM (SMF2000, Hitachi) were usedto examine the cross-sectional and three-dimensional images of theelectrodes after melt infiltration. Elemental distribution was analyzedby EDSmapping at identical positions. TEM (JEM-ARM200F, JEOL) andToF-SIMS (TOF.SIMS5-AD-GCIB, ION-TOF) were employed to char-acterize the electrode interface before and after electrochemicalmeasurements. ImageJ49 and Avizo (Thermo-Fisher Scientific) wereused to prepare two-dimensional and three-dimensional images,respectively.An X-ray diffractometer (D8 ADVANCE, Bruker) was used tocharacterize phases and lattice parameters in the solid electrolytes.Rietveld diffraction pattern analysis was performed using the crystal-lographic analysis software GSAS-II50.Electrochemical measurementsElectrochemical impedance spectroscopy (Solartron 1470E, Solar-tronAnalytical; E4990A, Keysight) was employed tomeasureNyquistand Bode plots, evaluating both the ionic conductivity of the solidelectrolyte in symmetric cells and interfacial resistivity in CR2016coin cells (full cells). The impedance spectra weremeasured using anAC voltage amplitude of 10mV under potentiostatic control, super-imposed on the open-circuit voltage, over a frequency range of7MHz to 0.1 Hz at various temperatures (298–333 K (±1 K)). DCmagnetron sputtering was employed to deposit 40 nm Pt films onboth sides of the composite solid electrolyte pellet for fabricatingsymmetric cells.A TOSCAT-3100 (Toyo system) connected to climate chamberwas used to perform charge–discharge tests on the ASSB coin cells.The initial charge–discharge tests were conducted at 298K ± 1 Kwith arate of 0.05 C in the voltage ranges of 0.5–2.5 V. Cycle performancetests were carried out at 333 K ± 1 K with a rate of 0.3 C. Rate perfor-mance tests were conducted at 298K with varying discharge ratesfrom 0.1C to 1 C.Data availabilityThe data that support the findings of this study are includedwithin thearticle, its Supplementary Information, Source Data, and Supplemen-tary Data files.The key calculation data generated in this study areprovided in the Supplementary Data as CIF files. Due to companypolicy, data deposition in public repositories is not possible. 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