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. 2026 Oct 5;16(46):60166–60178. doi: 10.1039/d6ra06621f

Enhancing ionic conductivity and interfacial stability of perovskite-type LLTO solid electrolytes via synergistic A-site Sr substitution and polymer layer engineering for safer lithium-ion batteries

Tausif Alam a, Lakshmi Hrushita Korlapati a,b, Sumit Khatua a,b, K Ramakrushna Achary c, L N Patro a,b,✉
PMCID: PMC13636766  PMID: 42836031

Abstract

The development of safe, high-performance solid electrolytes is critical for advancing next-generation solid-state lithium-ion batteries (SSLIBs). Perovskite-type Li0.33La0.56TiO3 (LLTO) has emerged as a promising solid electrolyte for SSLIBs owing to its high thermal and mechanical stability; however, its moderate ionic conductivity and interfacial instability with lithium metal hinder its practical application. In this work, the ionic conductivity of LLTO was enhanced through A-site doping via partial substitution with Sr using two different Sr precursors (SrCO3 and SrF2). Sr-doped LLTO solid electrolytes were synthesised using a solid-state reaction method. Structural and microstructural analyses confirmed the successful incorporation of Sr into LLTO without the formation of secondary phases. Ionic conductivity measurements revealed that Sr-doped LLTO synthesised using the SrF2 precursor (LLTO-SrF2) exhibited higher ionic conductivity than that synthesised using the SrCO3 precursor (LLTO-SrCO3). The enhanced conductivity of LLTO-SrF2 relative to LLTO-SrCO3 and pristine LLTO was attributed to differences in their lattice parameters, densities, and microstructures. LLTO-SrF2 exhibited a high room-temperature bulk ionic conductivity of 2.86 × 10−3 S cm−1. To further address the interfacial instability with lithium metal, a PEO-based protective layer (PL) was placed around the LLTO-SrF2 pellet ((LLTO-SrF2)PL) to suppress parasitic reactions between Ti4+ and lithium metal while improving interfacial contact. The symmetric cell (Li/(LLTO-SrF2)PL/Li) exhibited stable galvanostatic cycling for over 540 h with highly reversible and symmetric lithium plating/stripping behaviour. These results highlight the synergistic effect of Sr substitution in enhancing ionic conductivity and polymer layer engineering in improving electrochemical stability against lithium metal.

Development of a polymer-layer covered LLTO-SrF2 to overcome the interfacial instability with lithium metal.graphic file with name d6ra06621f-ga.webp

Introduction

Conventional organic liquid electrolytes employed in various energy-storage devices, including lithium-ion batteries (LIBs), are cost-effective and often exhibit high ionic conductivity. However, they are highly flammable and prone to leakage, leading to serious safety concerns.1–3 In contrast, solid-state LIBs (SSLIBs) employing solid electrolytes not only mitigate leakage and thermal runaway but also improve cycle life and energy density.4,5 The solid electrolytes widely investigated for their promising applications in SSLIBs mainly include oxide-based electrolytes such as NASICON-type Li1.5Al0.5Ge1.5(PO4)3 (LAGP), garnet-type Li7La3Zr2O12 (LLZO), perovskite-type Li0.33La0.56TiO3 (LLTO), and LISICON-type Li3.75Ge0.75P0.25O4 (LGPO).6–9 Each class of solid electrolyte possesses distinct advantages and limitations. Among them, garnet-type and NASICON-type solid electrolytes have received greater attention than perovskite-type solid electrolytes despite their comparable ionic conductivities. Although NASICON-type LAGP exhibits high ionic conductivity and superior electrochemical stability, its widespread application is limited by the high cost of GeO2, the Ge precursor. Garnet-type LLZO also exhibits high ionic conductivity; however, it is prone to instability upon exposure to moisture and CO2.10

Perovskite systems, in general, have attracted tremendous attention in materials science owing to their versatile crystal structures and remarkable electrical, magnetic, optical, and catalytic properties. They have found widespread applications in solar cells, LEDs, and ferroelectric devices, among others.11,12 However, compared with these applications, perovskite materials have been explored to a much lesser extent as solid electrolytes for energy-storage devices. Perovskite-type solid electrolytes adopt an ABO3 structure, where the smaller cation occupies the octahedrally coordinated B-site, while the larger cation resides at the 12-fold coordinated A-site.13 One of the major advantages of the perovskite structure is its structural flexibility, which enables substitution at both the A- and B-sites of the ABO3 framework without disrupting the overall crystal structure. This flexibility is governed by a broad tolerance factor range (0.75 < t < 1.0), allowing the incorporation of a wide variety of dopant ions with different ionic radii and oxidation states.14

Over the past decades, the Inline graphic family has attracted significant research interest owing to its diverse perovskite structures, which vary with lithium concentration. Among this family, Inline graphic with x = 0.11 (Li0.33La0.56TiO3, LLTO) offers several advantages, including high bulk Li+-ion conductivity on the order of 10−3 S cm−1 at room temperature and relatively simple synthesis.15 However, LLTO has not yet achieved commercial application owing to several critical limitations, including the reduction of Ti4+ to Ti3+ upon contact with lithium metal and its narrow electrochemical stability window.16 Furthermore, its ionic conductivity remains lower than that of conventional liquid electrolytes, which typically exhibit ionic conductivities on the order of 10−2 S cm−1 at room temperature.17 Various strategies have been explored to improve the ionic conductivity of solid electrolytes, including optimisation of synthesis methods, processing parameters such as calcination and sintering temperatures, doping, and crystal structure engineering. Furthermore, to mitigate the reduction of Ti4+ to Ti3+ in Ti-containing solid electrolytes, several approaches have been explored, including the use of protective interlayers and the introduction of small amounts of liquid electrolytes or ionic liquids at the solid electrolyte/Li interface.18–20

In this work, LLTO-based solid electrolytes were prepared using a simple and industry-friendly solid-state reaction method. Three main strategies were employed: (i) optimisation of the sintering temperature to achieve phase purity with enhanced density, (ii) doping to enhance ionic conductivity, and (iii) incorporation of a polymer-based protective layer at the solid electrolyte/Li interface to improve interfacial stability. The sintering temperature in the solid-state reaction method plays a vital role in determining the phase formation, density, and ionic conductivity of the solid electrolyte. Doping is a well-known strategy for tailoring the ionic conductivity of solid electrolytes. In the present study, strontium (Sr2+) was chosen as an effective dopant to partially substitute for La3+ site in LLTO. The larger ionic radius of Sr2+ compared with that of La3+ can modulate the ionic transport behaviour of LLTO by introducing defects associated with differences in ionic radius and charge. Based on promising results from previous studies, the Sr dopant level was fixed at 3 mol%.21 In previous studies, SrCO3 has commonly been used as the Sr source.22 In this study, both SrCO3 and SrF2 were employed as Sr precursors for the preparation of Sr-doped LLTO via the solid-state reaction method, with the latter additionally expected to enable fluorine substitution at the oxygen site. This study systematically investigates how the partial substitution of Sr2+ at the A-site (La3+) influences the crystal structure, ionic conductivity, electronic conductivity, and hydrophilicity of the LLTO electrolyte. Finally, a polymer-based protective layer was covered around the solid electrolyte pellet to prevent the reduction of Ti4+ to Ti3+ upon contact with lithium metal. Furthermore, its role at the LLTO/Li interface was evaluated in terms of lithium plating/stripping behaviour.

Experimental methods

Synthesis of bare and Sr-doped LLTO

Bare LLTO was synthesised via a solid-state reaction method using Li2CO3, La2O3, and TiO2 powders as starting materials. An additional 10 wt% Li2CO3 was added to compensate for lithium loss during calcination and sintering at elevated temperatures. Furthermore, Sr-doped LLTO materials were synthesised using the same method, employing Li2CO3, La2O3, and TiO2 powders along with either SrCO3 or SrF2 as the Sr precursor. The Sr dopant level was fixed at 3 mol% in both cases. Sr-doped LLTO prepared using SrCO3 and SrF2 as Sr precursors were denoted as LLTO-SrCO3 and LLTO-SrF2, respectively. The nominal compositions of LLTO, LLTO-SrCO3, and LLTO-SrF2 are Li0.33La0.56TiO3, Li0.36La0.53Sr0.03TiO3, and Li0.36La0.53Sr0.03TiO2.97F0.06, respectively. In the solid-state reaction method, the raw materials were initially mechanically milled in a zirconia jar for 6 h using ethanol as the milling medium in a planetary ball mill. The milled powder was dried at 80 °C for 6 h. The dried powder was then placed in an alumina crucible and calcined at 800 °C for 10 h in a muffle furnace. The calcined powder was subsequently reground and pressed into pellets (12 mm in diameter) under a uniaxial pressure of 300 MPa. The LLTO pellets were then sintered at different temperatures (1000, 1100, 1200, and 1250 °C) in air for 6 h to optimise their density and ionic conductivity. During sintering, the pellets were covered with mother powder to minimise lithium volatilisation and suppress the formation of impurity phases. The Sr-doped LLTO pellets were sintered at the optimum temperature of 1200 °C.

Preparation of polymer layer-covered LLTO-SrF2 pellet

The polymer–salt slurry was prepared using poly(ethylene oxide) (PEO) and lithium bis(trifluoromethanesulfonyl)imide (LiTFSI). PEO (80 wt%) was first dissolved in acetonitrile at room temperature and stirred for 6 h, followed by the addition of LiTFSI (20 wt%) to improve polymer–salt complexation, with further stirring for 4 h. The LLTO-SrF2 pellet was then immersed in the solution and allowed to dry slowly at room temperature to form a thin, uniform polymer layer around the pellet (Fig. 1). The resulting polymer layer uniformly covered the LLTO-SrF2 pellet surface and is hereafter denoted as (LLTO-SrF2)PL.

Fig. 1. Schematic representation showing the preparation of the polymer layer-covered LLTO-SrF2 pellet.

Fig. 1

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Characterisation

XRD measurements were carried out using a PANalytical Empyrean diffractometer with Cu Kα radiation to investigate the structural characteristics of bare LLTO and Sr-doped LLTO materials over a 2θ range of 10–90°. The crystallographic parameters were further refined using X'Pert HighScore software. Raman spectra were recorded using a portable Raman spectrometer (B&W Tek) equipped with a 532 nm excitation source. The hydrophilic properties were studied using a contact angle meter (Kyowa Interface Science Co. Ltd). Solid-state 7Li and 19F MAS NMR measurements were performed on a 400 MHz NMR spectrometer (JEOL, JNM-ECZL400R) equipped with a 3.2 mm MAS probe, operating at a spinning rate of 10 kHz. The surface morphology and elemental distribution of the pellets were examined using HR-SEM (Apreo S, Thermo Scientific). The densities of the pellets were determined using Archimedes' principle, with deionised (DI) water as the immersion medium. Ionic conductivity was measured by impedance spectroscopy using a Solartron 1260A impedance analyser over a frequency range of 1 Hz to 10 MHz. The dc polarisation method was employed to determine the ionic transport number using an ion-blocking configuration. The electrochemical stability window (ESW) was determined by LSV using the Metrohm Autolab. Furthermore, a symmetric lithium cell was fabricated to evaluate the interfacial stability between the solid electrolyte and lithium metal. Lithium plating/stripping measurements were conducted with a cycle duration of 20 min each. For the evaluation of the Li plating/stripping performance of the Li/LLTO/Li and Li/LLTO-SrF2/Li symmetric cells, a few drops of liquid electrolyte were applied at the Li/solid electrolyte interfaces. In contrast, no liquid electrolyte was used for the polymer-wrapped LLTO-SrF2 sample ((LLTO-SrF2)PL). In all cases, the thickness of the pellets was about 1 mm. In the present study, the uncoated cells with liquid electrolyte and polymer-coated cells without liquid electrolyte were assembled under identical conditions. In both cases, the cells were assembled in a Swagelok-type cell configuration inside an argon-filled glove box. Following assembly, the Swagelok cells were securely tightened using a spanner and properly sealed.

Results and discussion

XRD studies

The XRD patterns of bare LLTO pellets sintered at different temperatures (1000, 1100, 1200, and 1250 °C) are shown in Fig. 2(a).

Fig. 2. (a) XRD patterns of LLTO pellets sintered at different temperatures. (b) XRD patterns of LLTO and Sr-doped LLTO prepared using SrCO3 (LLTO-SrCO3) and SrF2 (LLTO-SrF2) as Sr precursors. (c) Diffraction peak shifts toward lower 2θ values upon Sr doping.

Fig. 2

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In all cases, the diffraction patterns are in good agreement with the standard tetragonal phase of LLTO (JCPDS 01-087-0935). However, unlike the pellets sintered at 1100, 1200, and 1250 °C, an additional peak is observed for the pellet sintered at 1000 °C, indicating an incomplete reaction at the lower sintering temperature. This additional peak is attributed to Li2TiO3 and is marked with a star symbol (*) in Fig. 2(a).

Fig. 2(b) shows the XRD patterns of Sr-doped LLTO pellets prepared using SrCO3 (LLTO-SrCO3) and SrF2 (LLTO-SrF2) as Sr precursors and sintered at 1200 °C. In both cases, the XRD patterns confirm that the tetragonal perovskite structure is retained upon Sr doping, with no detectable impurity phases. Furthermore, the diffraction peaks of the Sr-doped materials shift towards lower 2θ values, as shown in Fig. 2(c), indicating lattice expansion due to Sr incorporation into the LLTO lattice. This behaviour is mainly attributed to the larger ionic radius of the divalent Sr2+ (r = 1.44 Å) compared with that of trivalent La3+ (r = 1.36 Å). Overall, the XRD studies confirm the successful substitution of Sr at the La site following calcination at 800 °C and subsequent sintering at 1200 °C. The characteristic peak splitting further confirms the formation of the tetragonal phase in all samples (Fig. 2(c)).

Rietveld refinements were performed to determine the lattice parameters and unit cell volume. Fig. S1 shows the Rietveld-refined XRD patterns of LLTO, LLTO-SrCO3, and LLTO-SrF2 pellets sintered at 1200 °C. In all cases, the refinement results show good agreement between the observed and calculated patterns. The refined lattice parameters of bare LLTO are a = b = 3.8722 Å and c = 7.7424 Å, which are in good agreement with previous reports.21 The lattice parameters of the Sr-doped LLTO materials (LLTO-SrCO3 and LLTO-SrF2) are summarised in Table S1. As expected, the lattice parameters of the Sr-doped LLTO materials are slightly larger than those of bare LLTO.

Raman studies

The Raman spectra of LLTO sintered at different temperatures are shown in Fig. 3(a). In LLTO, the Raman bands originate from lattice distortions induced by Li/La ordering and Ti–O vibrations within the TiO6 octahedra.23,24 The number of bands and their relative positions are in good agreement with those predicted for the tetragonal structure of LLTO with the space group P4/mmm. In addition, an extra peak appears at approximately 866 cm−1 in the spectrum of LLTO sintered at 1000 °C. This additional band may be attributed to the vibrational mode of a lithium-titanium-based impurity phase. Notably, a corresponding Li2TiO3 impurity phase is also observed in the XRD pattern of the LLTO pellet sintered at 1000 °C, as shown in Fig. 2.

Fig. 3. (a) Raman spectra of LLTO sintered at different temperatures. (b) Raman spectra of LLTO-SrCO3 and LLTO-SrF2.

Fig. 3

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Fig. 3(b) shows the Raman spectra of LLTO-SrCO3 and LLTO-SrF2. In both cases, the characteristic bands are similar to those of bare LLTO, confirming the absence of impurity phases resulting from Sr doping in the LLTO lattice.

NMR studies

The NMR studies provide deep insights into Li+-ion dynamics in LLTO-based solid electrolytes. Fig. 4(a) shows the 7Li MAS NMR spectra of bare LLTO and Sr-doped LLTO (LLTO-SrCO3 and LLTO-SrF2). In all cases, the spectra exhibit a single, sharp, and intense central peak corresponding to the Inline graphic transition, indicating a single Li environment. In addition, two satellite bands arising from the Inline graphic and Inline graphic transitions, associated with the electric quadrupolar interactions, are observed.25,26Fig. 4(b) shows the 19F MAS NMR spectrum of LLTO-SrF2, which exhibits a single broad resonance centred at approximately −138.4 ppm, confirming the presence of fluorine in the sample.

Fig. 4. (a) 7Li MAS NMR spectra of bare LLTO and Sr-doped LLTO prepared using SrCO3 (LLTO-SrCO3) and SrF2 (LLTO-SrF2) as Sr precursors. (b) 19F MAS NMR spectrum of LLTO-SrF2.

Fig. 4

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Microstructure studies

Fig. 5 presents the top-view HR-SEM images of bare LLTO and Sr-doped LLTO (LLTO-SrF2). The HR-SEM image of LLTO-SrCO3 is provided in Fig. S2. The HR-SEM image of the bare LLTO pellet shows small, non-uniform grains with clearly visible grain boundaries. Furthermore, it indicates weak interconnection between neighbouring grains, which may hinder Li+-ion transport. A relatively compact surface morphology is observed in both Sr-doped LLTO pellets, indicating a reduced contribution of grain boundaries and improved Li+-ion migration compared with bare LLTO.

Fig. 5. HR-SEM images of (a) bare LLTO and (b) Sr-doped LLTO (LLTO-SrF2) pellets. Elemental mapping images of (c) bare LLTO and (d) LLTO-SrF2. EDX spectra of (e) bare LLTO and (f) LLTO-SrF2.

Fig. 5

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ImageJ software was used to calculate the average grain size of the bare LLTO, LLTO-SrCO3, and LLTO-SrF2 pellets. The average grain sizes of the bare LLTO, LLTO-SrCO3, and LLTO-SrF2 pellets are 1.8 µm, 2.2 µm, and 2.6 µm, respectively. A larger grain size often reduces the number of grain boundaries and, hence, the grain-boundary resistance, thereby facilitating smooth Li+-ion transport.27 Furthermore, the corresponding elemental mapping images and EDX spectra are presented in Fig. 5 and S2. The mapping images demonstrate the homogeneous distribution of the constituent elements, while the EDX analysis confirms the presence of the expected elements in both the bare and Sr-doped LLTO samples.

The density of the bare LLTO pellets sintered at different temperatures, along with those of the Sr-doped LLTO materials (LLTO-SrCO3 and LLTO-SrF2), was measured using Archimedes' principle with DI water as the immersion medium, and the obtained values are listed in Table S2. Among these, the LLTO-SrF2 pellet exhibits the highest relative density of 98.19%, which is attributed to improved grain contact and a larger grain size.

Hydrophilicity analysis

In the present study, contact angle measurements were performed using a water droplet to investigate the surface hydrophilicity of bare and Sr-doped LLTO materials. Bare LLTO exhibits a low average contact angle of 32.7°, indicating a highly hydrophilic surface. In contrast, LLTO-SrF2 exhibits a relatively high average contact angle of 45.5°, indicating reduced hydrophilicity compared with bare LLTO. Time-dependent contact angle measurements were performed over 20 s at 5 s intervals for all three systems, and the results are also included in Fig. 6. The rate of decrease in the contact angle was the fastest for LLTO (0.69° s−1), followed by LLTO-SrCO3 (0.45° s−1), and LLTO-SrF2 (0.44° s−1).

Fig. 6. (a) Contact angle images recorded at 5 s intervals and (b) the corresponding contact angle versus time graph.

Fig. 6

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Bare LLTO exhibits the fastest decrease in contact angle, indicating higher surface wettability, faster moisture adsorption, and lower moisture stability. In contrast, LLTO-SrCO3 and LLTO-SrF2 exhibit significantly slower decrease in contact angle, demonstrating weaker interactions with water molecules. The lowest rate of decrease in contact angle observed for LLTO-SrF2 suggests enhanced hydrophobicity and superior moisture resistance.

Ionic conductivity

Impedance spectroscopy was used to evaluate the conductivity of the solid electrolytes at both room and elevated temperatures. Fig. 7(a) presents the room-temperature Nyquist plots of LLTO sintered at different temperatures, while Fig. 7(b) shows the Nyquist plots of Sr-doped LLTO (LLTO-SrCO3 and LLTO-SrF2) along with LLTO sintered at 1200 °C. In all cases, the Nyquist plots exhibit two depressed semicircles in the high-frequency region and a spike in the low-frequency region. The intercept of the first semicircle with the real Z-axis corresponds to the bulk resistance, whereas the second semicircle is associated with the grain-boundary resistance. Fig. S3 shows the equivalent circuit employed for the analysis of the impedance data to determine the bulk and total resistances. Among the various sintering temperatures, LLTO sintered at 1200 °C exhibits the lowest bulk and grain-boundary resistances. Table 1 summarises the bulk and the total conductivities calculated from their respective resistance values using eqn (1).

graphic file with name d6ra06621f-t6.webp 1

where l is the electrolyte thickness, R is the resistance, and A is the electrode contact area. It is clearly seen from Table 1 that, among the various sintering temperatures investigated, LLTO sintered at 1200 °C exhibits the highest bulk (1.15 × 10−3 S cm−1) and total conductivity (3.04 × 10−6 S cm−1) at room temperature. The observed enhancement in conductivity is attributed to improved densification, enlarged grain size, and reduced grain-boundary resistance.

Fig. 7. (a) Room-temperature Nyquist plots of LLTO sintered at different temperatures. (b) Room-temperature Nyquist plots of LLTO-SrCO3 and LLTO-SrF2, along with LLTO sintered at 1200 °C.

Fig. 7

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Table 1. Room-temperature (RT) bulk and total conductivities of LLTO, LLTO-SrCO3, and LLTO-SrF2.

System Sintering temperature (°C) Bulk conductivity at RT (±0.10 S cm−1) Total conductivity at RT (±0.10 S cm−1) Activation energy (±0.01 eV)
1000 1.06 × 10−4 2.48 × 10−7 0.27
1100 4.87 × 10−4 6.21 × 10−7 0.24
LLTO 1200 1.15 × 10−3 3.04 × 10−6 0.28
1250 1.26 × 10−4 1.92 × 10−6 0.25
LLTO-SrCO3 1200 1.59 × 10−3 2.32 × 10−6 0.23
LLTO-SrF2 1200 2.86 × 10−3 7.85 × 10−6 0.25
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Fig. 7(b) shows the room-temperature Nyquist plots of LLTO sintered at 1200 °C, along with LLTO-SrCO3 and LLTO-SrF2. All three systems follow a similar trend, exhibiting two semicircles in the high-frequency region and a spike in the low-frequency region. Among these, LLTO-SrF2 shows the lowest bulk and total resistances. The obtained bulk and total conductivities of LLTO-SrF2 are 2.86 × 10−3 S cm−1, and 7.85 × 10−6 S cm−1, respectively, which are higher than those of bare LLTO and LLTO-SrCO3 (Table 1).

Fig. 8(a) presents the temperature-dependent bulk conductivity plots of LLTO sintered at different temperatures, whereas Fig. 8(b) shows the temperature-dependent bulk conductivity plots of LLTO-SrCO3 and LLTO-SrF2, along with LLTO sintered at 1200 °C. In all cases, the plots exhibit Arrhenius behaviour, Inline graphic, where σ0 is the pre-exponential factor.

Fig. 8. (a) Arrhenius plots of conductivity for LLTO sintered at different temperatures. (b) Arrhenius plots of conductivity for LLTO, LLTO-SrCO3, and LLTO-SrF2.

Fig. 8

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The activation energy (Ea) for each sample was determined from the slope of the temperature-dependent conductivity plots. The obtained activation energies are summarized in Table 1. Among these, LLTO-SrF2 exhibits the highest ionic conductivity with an activation energy of 0.25 eV. Sr-doped LLTO materials exhibited higher ionic conductivity than bare LLTO. This improvement can be attributed to the combined effects of Sr doping, increased grain size, and higher relative density. Among the Sr-doped LLTO materials prepared using SrCO3 and SrF2 as Sr precursors, LLTO-SrF2 exhibited the highest ionic conductivity. This superior conductivity may be associated with the synergistic incorporation of Sr and F into the LLTO structure, thereby modifying the defect chemistry and facilitating Li+ ion transport. Furthermore, a comparative table summarising the synthesis methods, relative densities, ionic conductivities, and activation energies of previously reported LLTO-based solid electrolytes and the solid electrolytes investigated in this work are provided in Table S3.

Electronic conductivity

An ideal solid electrolyte must be a purely ionic conductor because electronic conduction can cause electrical leakage or a short circuit in LIBs. The dc polarisation method was used to measure the electronic conductivity of LLTO, LLTO-SrCO3, and LLTO-SrF2 (Fig. 9). In this method, a small dc voltage (V) of 0.5 V was applied across an Ag/solid electrolyte/Ag cell, and the current was monitored as a function of time. The steady-state currents (Ie) of 0.021 µA, 0.014 µA, and 0.010 µA were obtained for LLTO, LLTO-SrCO3, and LLTO-SrF2, respectively.

Fig. 9. Polarisation current as a function of time.

Fig. 9

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The electronic conductivity can be calculated using Inline graphic. The calculated electronic conductivities of the LLTO, LLTO-SrCO3, and LLTO-SrF2 solid electrolytes are 6.84 × 10−9 S cm−1, 5.48 × 10−9 S cm−1, and 3.84 × 10−9 S cm−1, respectively, which are about three orders of magnitude lower than their ionic conductivities. These observations clearly demonstrate that the contribution of electronic conductivity to the total conductivity of the investigated perovskite electrolytes is negligible. Further, the ionic transport number (ti) was calculated using ti = 1 − (Ie/It), where It represents the initial current (total current). The calculated ti values for the LLTO, LLTO-SrCO3, and LLTO-SrF2 solid electrolytes are close to 1 (0.99), suggesting that ions are the dominant charge carriers. Among the two doped systems, LLTO-SrF2 was selected for further studies because of its higher ionic conductivity.

Electrochemical stability window

The electrochemical stability window (ESW) of the LLTO-SrF2 solid electrolyte was evaluated by linear sweep voltammetry (LSV) over a voltage range of 0.5–3.5 V at a scan rate of 0.1 mV s−1 using a Li/LLTO-SrF2/Au cell configuration. The LLTO-SrF2 solid electrolyte exhibited an ESW of up to 2.8 V vs. Li+/Li (Fig. S4).

Stability of LLTO and LLTO-SrF2 against lithium metal

Perovskite solid electrolytes suffer from electrochemical instability against lithium metal owing to their inherent redox activity. Upon direct contact with lithium metal, the transition-metal ions present in the perovskite framework (Ti ions in the present case) undergo reduction reactions, resulting in the structural and chemical degradation of the electrolyte.28,29 To evaluate the stability against lithium metal, lithium symmetric cells were assembled and used to investigate the long-term lithium plating/stripping behaviour of LLTO and LLTO-SrF2 pellets toward lithium metal. The configuration of the lithium symmetric cell employed for bare LLTO is schematically shown in Fig. 10(a). The symmetric cell was assembled inside an argon-filled glovebox by placing the LLTO pellet between two Li foil electrodes in a stacked configuration. A 5 µL cm−2 of liquid electrolyte (1 M LiPF6 in EC/DEC/DMC (v/v/v 1 : 1 : 1)) was applied to ensure better interfacial contact between the electrolyte and the electrodes. The stability of the symmetric cell was evaluated by performing Li plating/stripping cycles. The Li plating/stripping profiles of the symmetric cells, Li/LLTO/Li and Li/LLTO-SrF2/Li measured at a current density of 0.1 mA cm−2 are shown in Fig. 10(b) and S5, respectively. In both cases, the cells exhibited asymmetric cycling behaviour and uneven Li dissolution and deposition.

Fig. 10. (a) Schematic representation of the Li/LLTO/Li symmetric cell. (b) Li plating/stripping profile of the Li/LLTO/Li symmetric cell at a current density of 0.1 mA cm−2 at room temperature. (c) XRD patterns showing the formation of impurity phases in the LLTO pellets after different cycling durations of 2, 6, and 12 h, respectively.

Fig. 10

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Furthermore, the LLTO pellet was disassembled after Li plating/stripping cycling to examine changes in its crystal structure. It was observed that the colour of the LLTO pellet changed from dark yellow to black, as shown in Fig. 10(c), indicating structural modification upon contact with lithium metal. The XRD patterns of the LLTO pellet after 2, 6, and 12 h of cycling are displayed in Fig. 10(c). Various additional peaks were observed in the XRD patterns following Li plating/stripping cycling, as denoted by * TiO2, # La5Ti5O17, and & unknown phase, respectively. The XRD patterns clearly indicate that the number of additional peaks increased with increasing cycling duration. These additional peaks are associated with the reduction of Ti4+ ions present in the LLTO solid electrolyte upon contact with lithium metal, leading to the formation of Ti3+-containing phases.30 To rule out the chemical instability of the LLTO solid electrolyte toward the liquid electrolyte used, the LLTO pellet was immersed in the liquid electrolyte for 6 days. Fig. S6 shows the XRD patterns of bare LLTO and LLTO after immersion in the liquid electrolyte for 6 days. No noticeable changes were observed in the XRD patterns after immersion, indicating that the LLTO solid electrolyte remains chemically stable in the liquid electrolyte. Therefore, the structural degradation observed during cycling is attributed to the direct contact between LLTO and lithium metal. To suppress this interfacial reaction and improve compatibility with lithium metal, a protective polymer layer was designed (Fig. 1), and its role in preventing interfacial reactions between the solid electrolyte and lithium metal was investigated.

Effect of the protective polymer layer on the interfacial stability of LLTO-SrF2 against lithium metal

A few studies have demonstrated interfacial modification strategies to suppress parasitic reactions between Ti4+ in LLTO and lithium metal while improving interfacial contact. Lu et al. employed an electrolyte-soaked paper at the interface to mitigate the reaction between Li metal and the LLTO solid electrolyte (Li15/56Sr1/16La15/28TiO3).22 Yan et al. investigated the role of different dry- and gel-based polymer layers at the interface.31 In the present study, a protective polymer layer (80PEO-20LiTFSI) was covered around the LLTO-SrF2 pellet ((LLTO-SrF2)PL) to improve interfacial contact with lithium metal and prevent reactions between Ti4+ in LLTO-SrF2 and lithium metal. HRSEM measurements were performed to assess the morphology and elemental distribution of (LLTO)PL (Fig. S7). The HRSEM image reveals a densely packed morphology. The presence of La, Ti, O, F, S, N, and C was confirmed by EDS spectra, and their uniform elemental distribution was clearly observed in the respective elemental mapping images provided in Fig. S7. Fig. S8 shows the cross-sectional HRSEM images of (LLTO-SrF2)PL. The thickness of the polymer layer is about 50.1 µm. Furthermore, impedance measurements were carried out before coating (LLTO-SrF2) and after polymer layer coating ((LLTO-SrF2)PL) at room temperature (Fig. S9). As expected, the polymer layer coating introduces additional impedance, however, in this work, the focus is not on altering the conductivity of LLTO-SrF2 through polymer-layer coating, but rather on providing a stable, ionically conductive interfacial layer, improving contact with Li metal, and suppressing undesirable interfacial reactions.

The Li plating/stripping performance of the Li/(LLTO-SrF2)PL/Li symmetric cell at a current density of 0.1 mA cm−2 at room temperature is shown in Fig. 11(a). It was observed that the voltage profile decreased for up to 100 h, after which it became stable and smooth up to 540 h. In addition, all the Li dissolution and deposition curves exhibited a mirrored relationship, indicating excellent reversibility of the Li redox reaction in the cell and significantly improved stability for up to 540 h.

Fig. 11. (a) Li plating/stripping performance of the Li/(SrF2-LLTO)PL/Li symmetric cell at a current density of 0.1 mA cm−2 at room temperature. Enlarged views of the cycling performance: (b) 200 to 250 h (c) 400 to 450 h.

Fig. 11

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Fig. 11(b) and (c) shows the enlarged cycling voltage profiles from 200 to 250 h and 400 to 450 h, respectively. These results demonstrate the absence of any abrupt voltage drop and short-circuit development within the electrolyte, which can be attributed to the presence of the polymer protective layer between the LLTO-SrF2 solid electrolyte and the lithium electrodes. Furthermore, the polymer-covered solid electrolyte pellet was disassembled from the Li/(SrF2-LLTO)PL/Li symmetric cell after 540 h of Li plating/stripping cycling to examine changes in its crystal structure and evaluate the phase stability of the (SrF2-LLTO)PL electrolyte (Fig. 12). All the characteristic diffraction peaks of LLTO-SrF2 and the PEO-LiTFSI PL remained unchanged after cycling, with no observable peak shift or peak broadening. However, the formation of a thin or amorphous interphase cannot be ruled out. The absence of any additional diffraction peaks indicates that no secondary phases or decomposition products were formed during Li plating/stripping cycling for up to 540 h. This clearly confirms that the 80PEO-20LiTFSI protective polymer layer successfully mitigates the lithium-induced reduction of Ti4+ species and maintains a stable electrolyte/electrode interface.

Fig. 12. (a) XRD patterns of LLTO-SrF2 and (LLTO-SrF2)PL (b) XRD patterns showing the absence of impurity phases in (LLTO-SrF2)PL after 540 h of cycling.

Fig. 12

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Further, impedance spectroscopy measurements were performed before and after 540 h of Li plating/stripping. In both cases, the Nyquist plots exhibited similar behaviour, indicating that the repeated Li deposition/dissolution process does not significantly degrade the polymer-covered LLTO-SrF2 electrolyte, consistent with the XRD studies (Fig. S10).

Critical current density (CCD) analysis of (LLTO-SrF2)PL was performed using a stepwise galvanostatic Li plating/stripping test (Fig. 13). The current density was increased by 0.1 mA per cm2 per cycle, starting from 0.1 mA cm−2, with each Li plating/stripping step maintained for 1 h. The Li/(LLTO-SrF2)PL/Li cell exhibited stable and reversible voltage polarization with stable voltage profiles, demonstrating good tolerance toward repeated Li deposition and dissolution. The presence of the polymer layer improved the physical contact between the Li metal and the LLTO-SrF2 electrolyte. However, when the current density reached 0.9 mA cm−2, an abrupt change in the voltage response accompanied by a short circuit, was observed. Thus, the CCD of (LLTO-SrF2)PL was found to be 0.8 mA cm−2.

Fig. 13. Critical current density analysis of the Li/((LLTO-SrF2)PL)/Li symmetric cell.

Fig. 13

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Mechanism of the protective polymer layer at the Li/LLTO-SrF2 interface

Lithium metal, upon contact with LLTO-based electrolytes, undergoes oxidation according to the reaction: Li → Li+ + e−. Subsequently, the generated electrons are injected into the LLTO framework, where they reduce Ti4+ ions to Ti3+via the reaction: Ti4+ + e− → Ti3+.

This reduction process leads to the formation of an electronically conductive interphase within the LLTO electrolyte. The development of this conductive layer progressively diminishes the purely ionic nature of LLTO and promotes undesirable electronic transport through the electrolyte. Consequently, electron leakage increases, facilitating internal short-circuiting and ultimately degrading the electrochemical performance and stability of the cell. A schematic illustration of this mechanism is presented in Fig. 14(a).

Fig. 14. Schematic illustration of the proposed mechanism at the Li metal/LLTO interface: (a) without and (b) with a polymer protective layer.

Fig. 14

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In contrast, when a polymer protective layer is introduced between the lithium metal and the LLTO electrolyte, lithium oxidation still occurs during electrochemical operation. However, the polymer layer acts as an ionically conductive yet electronically insulating barrier. As a result, Li+ ions can readily migrate across the interface, while electron transport from the lithium metal to the LLTO electrolyte is effectively suppressed.31 The inhibition of electron transfer prevents the reduction of Ti4+ species, thereby maintaining the chemical and electrochemical stability of the LLTO structure. Consequently, the polymer protective layer enhances interfacial stability by preserving the ionic conduction characteristics of LLTO while blocking detrimental electronic conduction. A schematic representation of this mechanism is shown in Fig. 14(b).

Conclusions

A-Site substitution of LLTO with Sr was successfully carried out using SrCO3 and SrF2 as Sr precursors via a solid-state synthesis method. Structural characterization confirmed the formation of phase-pure perovskite materials exhibiting a tetragonal structure. Among the studied compositions, Sr-doped LLTO prepared using SrF2 as the Sr precursor (LLTO-SrF2) exhibited the highest bulk ionic conductivity of 2.86 × 10−3 S cm−1 at room temperature, with an activation energy of 0.25 eV and an ion transport number close to unity. The enhanced ionic conductivity of LLTO-SrF2 could be attributed to the combined effects of Sr incorporation, possible F substitution at the O sites, increased relative density, and improved microstructural characteristics. Furthermore, to improve the interfacial stability and mitigate the reduction of Ti4+ upon contact with lithium metal, a protective polymer layer (80PEO-20LiTFSI) was covered around the LLTO-SrF2 pellet ((LLTO-SrF2)PL). The symmetric cell using (LLTO-SrF2)PL demonstrated stable cycling for over 540 h with highly reversible lithium plating/stripping behaviour, confirming excellent electrochemical stability. Overall, this study demonstrates that the combined strategy of Sr substitution and interfacial engineering through the covering of a protective polymer layer is highly effective in improving both bulk ionic transport and interfacial compatibility. These findings provide a viable pathway for the development of high-performance, safe perovskite-based solid electrolytes for next-generation SSLIBs.

Author contributions

Tausif Alam: investigation, methodology, formal analysis, writing – original draft. Lakshmi Hrushita Korlapati: methodology. Sumit Khatua: methodology, writing – review & editing. K. Ramakrushna Achary: methodology. L. N. Patro: conceptualization, writing – review & editing, project administration, supervision.

Conflicts of interest

The authors declare no competing financial interest.

Supplementary Material

RA-016-D6RA06621F-s001

Acknowledgments

The author L. N. Patro acknowledges SRM Trust AP (SRMAP/URG/E&PP/2022-23/001) for its support and thanks the Department of Science and Technology, Delhi, India for supporting this research under the DST-FIST (SR-FST-CS-I-2021-219(C)) by providing the NMR facility.

Data availability

Any data that support the findings of this study are included within the article, and all the data used in the manuscript are available.

Supplementary information (SI) is available. See DOI: https://doi.org/10.1039/d6ra06621f.

References

  1. Chen F. Yang D. Zha W. Zhu B. Zhang Y. Li J. Gu Y. Shen Q. Zhang L. Sadoway D. R. Solid polymer electrolytes incorporating cubic Li7La3Zr2O12 for all-solid-state lithium rechargeable batteries. Electrochim. Acta. 2017;258:1106–1114. doi: 10.1016/j.electacta.2017.11.164. [DOI] [Google Scholar]
  2. Cao J. Meng J. Shen H. Ye X. Liu H. Guo C. Li J. Zhang G. Bao W. Yu F. Solid-state electrolytes for next-generation batteries: recent advances and interfacial challenges. J. Power Sources. 2025;654:237870. doi: 10.1016/j.jpowsour.2025.237870. [DOI] [Google Scholar]
  3. Zha W. Xu Y. Chen F. Shen Q. Zhang L. Cathode/electrolyte interface engineering via wet coating and hot pressing for all-solid-state lithium battery. Solid State Ionics. 2019;330:54–59. doi: 10.1016/j.ssi.2018.12.008. [DOI] [Google Scholar]
  4. Huo H. Luo J. Thangadurai V. Guo X. Nan C. W. Sun X. Li2CO3: a critical issue for developing solid garnet batteries. ACS Energy Lett. 2020;5:252–262. doi: 10.1021/acsenergylett.9b02401. [DOI] [Google Scholar]
  5. Zhang J. Zhao N. Zhang M. Li Y. Chu P. K. Guo X. Di Z. Wang X. Li H. Flexible and ion-conducting membrane electrolytes for solid-state lithium batteries: dispersion of garnet nanoparticles in insulating polyethylene oxide. Nano Energy. 2016;28:447–454. doi: 10.1016/j.nanoen.2016.09.002. [DOI] [Google Scholar]
  6. Okumura T. Taminato S. Miyazaki Y. Kitamura M. Saito T. Takeuchi T. Kobayashi H. LISICON-based amorphous oxide for bulk-type all-solid-state lithium-ion battery. ACS Appl. Energy Mater. 2020;3:3220–3229. doi: 10.1021/acsaem.9b01949. [DOI] [Google Scholar]
  7. Lu J. Li Y. Perovskite-type Li-ion solid electrolytes: a review. J. Mater. Sci.: Mater. Electron. 2021;32:9736–9754. doi: 10.1007/s10854-021-05699-8. [DOI] [Google Scholar]
  8. Khatua S. Achary K. R. Sasikumar K. Korlapati L. H. Patro L. N. Engineering dense superionic Li1+xAlxTi2-x(PO4)3 solid electrolytes for safer solid-state Li-ion batteries: Impact of sintering temperature and Al3+ doping. Solid State Ionics. 2026;436:117128. doi: 10.1016/j.ssi.2026.117128. [DOI] [Google Scholar]
  9. Alam T. Das A. Review on impurity and conductivity issues of garnet type Li7La3Zr2O12: mechanisms, solutions, and perspectives. Energy Fuels. 2023;37:15267–15282. doi: 10.1021/acs.energyfuels.3c01821. [DOI] [Google Scholar]
  10. Sharafi A. Kazyak E. Davis A. L. Yu S. Thompson T. Siegel D. J. Dasgupta N. P. Sakamoto J. Surface chemistry mechanism of ultra-low interfacial resistance in the solid-state electrolyte Li7La3Zr2O12. Chem. Mater. 2017;29:7961–7968. doi: 10.1021/acs.chemmater.7b03002. [DOI] [Google Scholar]
  11. Hu X. Cheng X. Qin S. Yan G. Malzbender J. Qiang W. Huang B. Mechanical and electrochemical properties of cubic and tetragonal LixLa0.557TiO3 perovskite oxide electrolytes. Ceram. Int. 2018;44:1902–1908. doi: 10.1016/j.ceramint.2017.10.129. [DOI] [Google Scholar]
  12. Zhang L. Mei L. Wang K. Lv Y. Zhang S. Lian Y. Liu X. Ma Z. Xiao G. Liu Q. Zhai S. Zhang S. Liu G. Yuan L. Guo B. Chen Z. Wei K. Liu A. Yue S. Niu G. Pan X. Sun J. Hua Y. Wu W. Q. Di D. Zhao B. Tian J. Wang Z. Yang Y. Chu L. Yuan M. Zeng H. Yip H. L. Yan K. Xu W. Zhu L. Zhang W. Xing G. Gao F. Ding L. Advances in the application of perovskite materials. Nano-Micro Lett. 2023;15:177. doi: 10.1007/s40820-023-01140-3. [DOI] [PMC free article] [PubMed] [Google Scholar]
  13. Lu J. Li Y. Kong Y. Zhang N. Study of structure and conductivity of Li3/8Sr7/16-3x/2LaxZr1/4Nb3/4O3 solid electrolytes. Ceram. Int. 2018;44:4744–4750. doi: 10.1016/j.ceramint.2017.12.058. [DOI] [Google Scholar]
  14. Stramare S. Thangadurai V. Weppner W. Lithium lanthanum titanates: a review. Chem. Mater. 2003;15:3974–3990. doi: 10.1021/cm0300516. [DOI] [Google Scholar]
  15. Yu K. Tian Y. Gu R. Jin L. Ma R. Sun H. Xu Y. Xu Z. Wei X. Ionic conduction, colossal permittivity and dielectric relaxation behavior of solid electrolyte Li3xLa2/3-xTiO3 ceramics. J. Eur. Ceram. Soc. 2018;38:4483–4487. doi: 10.1016/j.jeurceramsoc.2018.05.023. [DOI] [Google Scholar]
  16. Chen C. H. Amine K. Ionic conductivity, lithium insertion and extraction of lanthanum lithium titanate. Solid State Ionics. 2001;144:51–57. doi: 10.1016/S0167-2738(01)00884-0. [DOI] [Google Scholar]
  17. Karuppasamy K. Theerthagiri J. Vikraman D. Yim C. J. Hussain S. Sharma R. Maiyalagan T. Qin J. Kim H. S. Ionic liquid-based electrolytes for energy storage devices: a brief review on their limits and applications. Polymers. 2020;12:918. doi: 10.3390/polym12040918. https://dx.doi.org/10.3390/POLYM12040918 [DOI] [PMC free article] [PubMed] [Google Scholar]
  18. Liu Y. Sun Q. Zhao Y. Wang B. Kaghazchi P. Adair K. R. Li R. Zhang C. Liu J. Kuo L. Y. Hu Y. Sham T. K. Zhang L. Yang R. Lu S. Song X. Sun X. Stabilizing the interface of NASICON solid electrolyte against Li metal with atomic layer deposition. ACS Appl. Mater. Interfaces. 2018;10:31240–31248. doi: 10.1021/acsami.8b06366. [DOI] [PubMed] [Google Scholar]
  19. Xia Q. Yuan S. Zhang Q. Huang C. Liu J. Jin H. Designing the interface layer of solid electrolytes for all-solid-state lithium batteries. Adv. Sci. 2024;11:e2401453. doi: 10.1002/advs.202401453. [DOI] [PMC free article] [PubMed] [Google Scholar]
  20. Tang J. Wang L. Tian C. Chen C. Huang T. Zeng L. Yu A. Double-protected layers with solid–liquid hybrid electrolytes for long-cycle-life lithium batteries. ACS Appl. Mater. Interfaces. 2022;14:4170–4178. doi: 10.1021/acsami.1c21457. [DOI] [PubMed] [Google Scholar]
  21. Zhang S. Zhao H. Guo J. Du Z. Wang J. Świerczek K. Characterization of Sr-doped lithium lanthanum titanate with improved transport properties. Solid State Ionics. 2019;336:39–46. doi: 10.1016/j.ssi.2019.03.015. [DOI] [Google Scholar]
  22. Lu J. Li Y. Ding Y. Structure, stability, and ionic conductivity of perovskite Li2x-ySr1-x-yLayTiO3 solid electrolytes. Ceram. Int. 2020;46:7741–7747. doi: 10.1016/j.ceramint.2019.11.277. [DOI] [Google Scholar]
  23. Rana J. Bilash M. R. I. Islam M. K. Rahman M. L. Optimization of Sr and Zr co-doping in lithium lanthanum titanate perovskite electrolytes for high-performance solid-state battery applications. J. Mater. Sci.: Mater. Electron. 2025;36:1616. doi: 10.1007/s10854-025-15697-9. [DOI] [Google Scholar]
  24. Le H. T. T. Kalubarme R. S. Ngo D. T. Jang S. Y. Jung K. N. Shin K. H. Park C. J. Citrate gel synthesis of aluminum-doped lithium lanthanum titanate solid electrolyte for application in organic-type lithium-oxygen batteries. J. Power Sources. 2015;274:1188–1199. doi: 10.1016/j.jpowsour.2014.10.146. [DOI] [Google Scholar]
  25. Hasegawa G. Kuwata N. Hashi K. Tanaka Y. Takada K. Lithium-ion diffusion in perovskite-type solid electrolyte lithium lanthanum titanate revealed by pulsed-field gradient nuclear magnetic resonance. Chem. Mater. 2023;35:3815–3824. doi: 10.1021/acs.chemmater.2c03340. [DOI] [Google Scholar]
  26. Emery J. Buzare J. Y. Bohnke O. Fourquet J. L. Lithium-7 NMR and ionic conductivity studies of lanthanum lithium titanate electrolytes. Solid State Ionics. 1997;99:41–51. doi: 10.1016/S0167-2738(97)00202-6. https://dx.doi.org/10.1016/s0167-2738(97)00202-6 [DOI] [Google Scholar]
  27. Khatua S. Rao Y. B. Achary K. R. Patro L. N. Li-ion transport studies of NASICON-type LiZr2(PO4)3 solid elctrolyte crystallizing in rhombohedral structure at room tempearature. Surf. Interfaces. 2023;41:103212. doi: 10.1016/j.surfin.2023.103212. [DOI] [Google Scholar]
  28. Zheng Z. Fang H. Liu Z. Wang Y. A fundamental stability study for amorphous LiLaTiO3 solid electrolyte. J. Electrochem. Soc. 2015;162:A244–A248. doi: 10.1149/2.0011503jes. [DOI] [Google Scholar]
  29. Ring J. Laa L. Limbeck A. Vonk V. Volkov S. Nenning A. Fleig J. Electrochemical stability window and electrolyte breakdown mechanisms of lithium lanthanum titanate. J. Electrochem. Soc. 2023;170:060509. doi: 10.1149/1945-7111/acd818. [DOI] [Google Scholar]
  30. Yang K. Y. Leu I. C. Fung K. Z. Hon M. H. Hsu M. C. Hsiao Y. J. Wang M. C. Mechanism of the interfacial reaction between cation-deficient La0.56Li0.33TiO3 and metallic lithium at room temperature. J. Mater. Res. 2008;23:1813–1825. doi: 10.1557/JMR.2008.0255. https://dx.doi.org/10.1557/jmr.2008.0255 [DOI] [Google Scholar]
  31. Yan S. Al-Salih H. Yim C. H. Merati A. Baranova E. A. Weck A. Abu-Lebdeh Y. Engineered interfaces between perovskite La2/3xLi3xTiO3 electrolyte and Li metal for solid-state batteries. Front. Chem. 2022;10:966274. doi: 10.3389/fchem.2022.966274. [DOI] [PMC free article] [PubMed] [Google Scholar]

Associated Data

This section collects any data citations, data availability statements, or supplementary materials included in this article.

Supplementary Materials

RA-016-D6RA06621F-s001
RA-016-D6RA06621F-s001.pdf (1.8MB, pdf)

Data Availability Statement

Any data that support the findings of this study are included within the article, and all the data used in the manuscript are available.

Supplementary information (SI) is available. See DOI: https://doi.org/10.1039/d6ra06621f.

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