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## Creator

[Yuki Nomura](https://orcid.org/0000-0002-6091-5902), Kazuo Yamamoto, [Naoaki Kuwata](https://orcid.org/0000-0002-0736-6967), Tsukasa Hirayama

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This document is the Accepted Manuscript version of a Published Work that appeared in final form in ACS Energy Letters, copyright © 2025 American Chemical Society after peer review and technical editing by the publisher. To access the final edited and published work see https://doi.org/10.1021/acsenergylett.5c00209.[In Copyright](http://rightsstatements.org/vocab/InC/1.0/)

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[Imaging Phase Boundary Kinetics in Lithium Titanate Using Operando Electron Energy-Loss Spectroscopy](https://mdr.nims.go.jp/datasets/97c23adc-85c7-46a3-a5da-4b1ec0118a4a)

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Template for Electronic Submission to ACS JournalsImaging the temperature-dependent ionic diffusion in lithium titanate using operando electron energy-loss spectroscopyYuki Nomura,1,* Kazuo Yamamoto,1 Naoaki Kuwata 2 and Tsukasa Hirayama 11Nanostructures Research Laboratory, Japan Fine Ceramics Center, 2–4–1 Mutsuno, Atsuta-ku, Nagoya, Aichi, 456–8587, Japan.2Research Center for Energy and Environmental Materials, National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan.* To whom correspondence should be addressed. E-mail: y_nomura@jfcc.or.jpLithium titanate accommodates and releases lithium ions through a phase separation. The dynamics of the phase boundary movement are critical to battery performance, particularly for maximizing the charge and discharge rates. However, the processes that facilitate ionic movement in lithium titanate are not well understood. Here, we visualize the phase boundary movement through the dynamic observation of Li distribution using operando scanning transmission electron microscopy coupled with electron energy-loss spectroscopy. The temperature dependent ionic diffusion during charging and discharging was evaluated with a maximum temporal resolution of 1.5 s/image. The rate constants of the phase boundary movement were determined to be 4.0 ± 0.7 × 10−11 cm2/s at 30 °C and 1.9 ± 0.6 × 10−9 cm2/s at 105 °C during Li extraction and 3.6 ± 0.9 × 10−13 cm2/s at 30 °C and 3.2 ± 0.3 × 10−11 cm2/s at 105 °C during Li insertion. The activation energy for Li-ion diffusion was calculated to be 0.49 eV for Li4Ti5O12 and 0.59 eV for Li7Ti5O12. The relatively low activation energy of 0.49 eV is the reason why lithium titanate exhibits high-rate discharge performance.TOC GRAPHICSLithium titanate (Li4+xTi5O12: LTO) has drawn significant interest as an anode material for solid-state Li-ion batteries because of its near-zero volume change and high redox potential at 1.55 V vs. Li during charging and discharging.1–7 These characteristics effectively mitigate the chemo-mechanical degradation at the interface with the solid electrolyte.8–12 LTO accommodates and releases Li ions through a two-phase reaction between an initial disordered spinel phase (Li4Ti5O12: ) and a rock-salt phase (Li7Ti5O12: ).2,13 Because the phase transition is accompanied by the movement of a phase boundary, the kinetics of the phase boundary movement are critical to battery performance, particularly for maximizing the charge/discharge rate.14,15 However, because ionic transport is an inherently nonequilibrium process, the specific mechanisms that enable facile ionic transport in LTO are poorly understood. For example, Li-ion diffusivities are reported to be low in the end-member phases (Li4Ti5O12 and Li7Ti5O12), which conflicts with the high-rate capability observed in electrochemical reactions.16–21 Therefore, investigating the Li-ion diffusion kinetics within the two phases in greater detail is important for the development of improved LTO anode materials and their analogues. Among available characterization techniques, electrochemical analyses, such as chronoamperometry, galvanostatic/potentiostatic intermittent titration technique (GITT/PITT), and electrochemical impedance spectroscopy (EIS), are widely used for the analysis of ionic diffusion and phase boundary movement.2,22,23 Nuclear magnetic resonance spectroscopy (NMR) and secondary ion mass spectrometry (SIMS) with isotope exchange methods are also used for the analysis of ionic diffusion.24–27 For two-phase reaction materials, potential-step chronoamperometry, GITT, and PITT have been used to investigate the phase transition in graphite and LiFePO4.28–34 In these studies, the rate constant of the phase boundary movement was evaluated from the current or voltage response. However, the values calculated using these electrochemical analyses are averages over the entire cell and depend on several assumptions, including uniform diffusion length and electrochemical reaction, and may not hold true for practical cells. Moreover, these approaches cannot be used to study the dynamics in local regions at the nanometer scale, including surfaces, defects, and grain boundaries. To measure the moving phase boundary directly, in situ or operando nanoscale microscopy is promising.35–41 In particular, operando scanning transmission electron microscopy (STEM) coupled with electron energy-loss spectroscopy (EELS) is one of the most promising analytical tools for this purpose because of the quantitative detectability of Li ions at the nanometer scale.42–44 Recently, our group has directly visualized Li-ion diffusion and phase transition in LTO using operando STEM-EELS.45 A spatially asymmetric phase separation during Li insertion and extraction was observed. However, the temporal resolution of the analysis was about 5 min/image and was only conducted at room temperature. These limitations make it challenging to analyze the kinetics of the phase boundary movement in greater detail. Therefore, achieving operando STEM-EELS with much higher temporal resolution at various temperatures is essential for comprehending Li-ion diffusion kinetics in LTO. In this study, we present operando STEM-EELS with high temporal resolution and precise control of sample temperature. We investigated the temperature dependence of the phase boundary movement caused by electrochemical Li insertion and extraction. Our temperature-resolved operando STEM-EELS visualizes the dynamic changes in the nanoscale Li-ion distribution with a maximum temporal resolution of 1.5 s/image. This high spatial and temporal resolution enables the real-time monitoring of Li-ion movement in a battery. In addition, we determined the parabolic rate constants of the phase boundary movement at various temperatures from the observed Li-ion maps. Finally, we evaluated the activation energy of Li-ion diffusion from the Arrhenius-type temperature dependence of the parabolic rate constant.Fig. 1. Solid-state Li-ion battery. (a) Schematic of the solid-state Li-ion battery used in this study. (b) ADF-STEM image around an LTO particle. (c) Typical charge and discharge curves of the battery cycled at room temperature in a transmission electron microscope. Reproduced from Ref. (45) with permission from the Royal Society of Chemistry. (d) Nyquist diagram of impedance spectra at different temperatures. (e) Chronoamperograms at 1.6 V (Li extraction) and 0.4 V (Li insertion) at different temperatures.Figure 1a shows a schematic of a solid-state Li-ion battery cell used in this study. LTO particles were embedded on the surface of a 70Li2S-30P2S5 (LPS) glass ceramic solid electrolyte layer. An In-Li alloy film was used as the counter electrode. A Pt film was deposited on the LTO particles by sputtering. Details of cell fabrication are shown in the supplementary information of our previous study.45 The area marked by a dashed rectangle in Fig. 1a was thinned using a focused Ga-ion beam (FIB) for operando STEM-EELS observations. The thinned battery cell was mounted on a vacuum-transfer TEM holder with two biasing electrodes and a heating stage (Mel-build Co.). An annular dark-field (ADF) STEM image of the thinned area is shown in Fig. 1b. A few micron-sized LTO particles were connected to the LPS particles and the Pt film, yielding ionic and electronic conduction pathways for the electrochemical reaction. Figure 1c shows the typical charge–discharge curves of the cell operated at room temperature in a transmission electron microscope. A potential plateau at 0.95 V vs. In-Li was observed. The asymmetric charge–discharge voltage profiles46 are a result of microscale asymmetric phase distributions and different ionic diffusivities of the end-member phases and surfaces.45 Figure 1d shows EIS spectra of the thinned battery cell at 30, 45, 60, 75, 90, and 105 °C obtained in a transmission electron microscope. The semicircular arcs in the high-frequency region are assigned to the resistance of the LPS layer. The slopes observed in the lower-frequency region correspond to the Warburg impedance, which is related to Li-ion diffusion. The diameters of the semicircular arcs were dependent on the cell temperature, showing that the ionic conductivity of LPS increased with the increase in temperature. After fully inserting Li at 0.4 V vs. In-Li, we performed potential-step chronoamperometry at 1.6 V vs. In-Li for Li extraction; once extraction was complete, we repeated the procedure at 0.4 V vs. In-Li to insert Li ions. Figure 1e shows the potential-step chronoamperograms obtained at different temperatures. The applied voltage of 1.6 and 0.4 V vs. In-Li is much higher and lower than the redox potential of LTO (0.95 V vs. In-Li). The large polarization at the interface between LTO and LPS results in a fast electrochemical reaction at the interface, and thus, Li-ion diffusion in LTO mainly limited the electrochemical reaction and the current. Higher temperatures induce faster Li-ion diffusion and movement of the phase boundary in LTO, resulting in a higher current at the beginning of the measurement. As the electrochemical reaction proceeds rapidly at the interface, mass transfer near the interface becomes the limiting factor, causing the current to decay. We observed this behavior in real space using operando STEM-EELS.Fig. 2. Change in the Li distribution in LTO during Li extraction visualized by operando STEM-EELS. (a) ADF-STEM image of the scanned area for STEM-EELS. (b) Crystal orientation map along the direction of the electron beam. (c) Change in the Li distribution during chronoamperometry at 1.6 V at 105, 90, 75, 60, 45, and 30 °C. Dashed yellow lines show the phase boundary between Li-poor Li4Ti5O12 and Li-rich Li7Ti5O12 phases. The timestamps on the images show the time elapsed since the start of the reaction. Figure 2a shows an ADF-STEM image of the scanned area for operando STEM-EELS observation. The LTO and LPS particles are connected at the area indicated by the dashed yellow lines. Li ions are extracted in the direction shown by the yellow arrows. Figure 2b shows a crystal orientation map of the LTO particle along the direction of the electron beam, which was obtained by crystal orientation mapping.47,48 The large LTO particle in the center of the image had a nearly uniform crystal orientation. As shown, there is a 2.9° misorientation along the line marked by the black arrows, probably arising from the small-angle grain boundary. However, this misorientation did not significantly change the ionic diffusion because the phase boundary movement was not affected in this region (see Fig. 2c). Figure 2c shows dynamic changes in the Li-ion distribution during the chronoamperometry of 1.6 V (Li extraction) at 30, 45, 60, 75, 90, and 105 °C. The temporal resolution of the Li-ion maps was 1.5 s/image. In these images, the bright regions represent the Li-rich Li7Ti5O12 phase, whereas the dark regions represent the Li-poor Li4Ti5O12 phase.49 Crucially, the phase boundary movement could be clearly seen. Because the reaction proceeds from the solid electrolyte side rather than from the current collector side, it indicates that ionic diffusion is the rate-limiting factor. Note that because the lattice constant difference between the Li4Ti5O12 and Li7Ti5O12 phases is less than 1%, it is difficult to distinguish the two phases using electron diffraction (Fig. S1) or ADF-STEM images (Fig. S2). The core-loss EELS analysis of the Ti-L and O-K edges are shown in Fig. S2. The results confirm that in Li4Ti5​O12 regions, the Ti-L edge is indicative of Ti4+, whereas in Li7Ti5​O12 regions, the Ti-L edge exhibits partially reduced features consistent with a mixture of Ti3+ and Ti4+. These observations are fully in line with the expected phase changes during Li extraction.13,49 Panels in Fig. 2c shows Li-ion distributions captured at times when the phase boundaries were approximately in the same position at different temperatures (see the Supplementary Materials for videos of this movement). At 105 °C, Li ions were rapidly extracted from the interface marked by the yellow dashed lines shown in Fig. 2a. In just 6 s, half of the Li7Ti5O12 phase transformed into the Li4Ti5O12 phase, and almost the entire region transformed into the Li4Ti5O12 phase within 15 s. At 90, 75, 60, 45, and 30 °C, the changes in the Li distribution were similar to those observed at 105 °C, except for the required time. This indicates that the limiting factor of the electrochemical reaction was the same, that is, the ionic diffusion in LTO. The times required to extract almost all Li ions from the LTO particle were 25.5, 46.5, 91.5, 181.5, and 387 s at 90, 75, 60, 45, and 30 °C, respectively. Note that the total electron dose in this experiment was approximately 5 × 103 electron/Å2, which is 100-times lower than that used for atomic resolution STEM observations.50 Therefore, we believe that electron irradiation damage does not significantly impact the results. A detailed quantitative discussion will be provided later. To confirm the reproducibility and representativeness of our measurements, we conducted the same experiments on 4 particles under the same conditions (Fig. S3 and S4), all of which yielded results consistent with those shown in Fig. 2. A more detailed, quantitative analysis of the phase boundary movement across these particles will be presented later. Furthermore, to analyze the changes in the Li diffusion behavior under diffusion-limited and reaction-limited conditions, we compared two scenarios: a constant voltage (CV) of 1.6 V and a constant current (CC) of 2 µA. Figure S5 shows the comparison for the same particle. Under the CV condition (Figures S5b–c, diffusion-limited case), the reaction proceeded from two interfaces within the field of view, whereas under the CC condition (Figures S5d–e, reaction-limited case), the reaction proceeded from only one interface. The difference in reaction distribution reflect a change in the rate-limiting step, paralleling the distinction between "shrinking core" and "intercalation wave" processes commonly reported in liquid-based LIB systems.51Fig. 3. Change in the Li distribution in LTO during Li insertion visualized by operando STEM-EELS. (a) ADF-STEM image of the scanned area for STEM-EELS. (b) Crystal orientation map along the direction of the electron beam. (c) Change in the Li distribution during chronoamperometry at 0.4 V at 105, 90, 75, 60, 45, and 30 °C. Dashed yellow lines show the phase boundary between Li-poor Li4Ti5O12 and Li-rich Li7Ti5O12 phases. The timestamps on the images show the time elapsed since the start of the reaction. Figure 3 shows the dynamic changes in the Li-ion distribution during the chronoamperometry of 0.4 V (Li insertion) at 30, 45, 60, 75, 90, and 105 °C. Figure 3a and 3b shows an ADF-STEM image and the crystal orientation map of the scanned area for operando STEM-EELS. The dashed yellow line outlines the interface. Figure 3c shows dynamic changes in the Li-ion distribution during Li insertion at various temperatures. Dashed yellow lines show the phase boundary between the Li-poor and Li- rich phases. During Li insertion, a core–shell structure composed of a Li-poor core and a Li-rich shell was observed, even though the LTO particle was not fully covered by the electrolyte (see the rightmost panel at 75 °C). This core–shell formation is attributed to differences in intrinsic Li-ion diffusivities among Li7​Ti5​O12 (slower), surface (intermediate), and Li4Ti5​O12 (faster), as well as to the partial coverage of the LTO particle with LPS in the solid-state battery configuration. During Li insertion, the influence of surface diffusion becomes more pronounced, resulting in the formation of a core–shell structure. A detailed explanation of the core–shell formation mechanism is described in our previous study45 and Fig. S6. Surface diffusion also occurred on the thinned surfaces oriented perpendicular to the electron beam during Li insertion. Figure 3d presents a 3D schematic of the lithiated lamella. A 3D Li7Ti5O12 shell was formed, and in the projection along the electron beam, the outer Li7Ti5O12 shell overlapped with the inner Li4Ti5O12​ core, leading to intermediate Li concentrations between these two phases. This explains why the two-phase boundary is less visible during Li insertion than that during Li extraction. Figure 3c illustrates that Li insertion takes substantially longer compared to Li extraction. At 30 °C, even after approximately 2 h, Li insertion remains incomplete. To verify the reproducibility and representativeness of our findings, we repeated these experiments on 3 particles under identical conditions (Fig. S7 and S8), and obtained results consistent with those in Fig. 3. Wagner investigated a diffusion-limited phase transformations with a moving phase boundary.29,52,53 Based on this method, when Li-ion diffusion is the rate-determining step, the position of the phase boundary (𝜉) is correlated with time (t), as shown by equation (1):where k is the parabolic rate constant of the phase boundary movement. When a new phase nucleates from an original phase, the relationship between the chemical diffusion coefficient of the original phase (DⅠ), chemical diffusion coefficient of the growing phase (DⅡ), and k is described by equation (2):where 𝛶 is a dimensionless parameter. This equation shows that the phase boundary movement is mainly affected by the diffusivity of Li ions in the growing phase, which was Li4Ti5O12 during Li extraction and Li7Ti5O12 during Li insertion. Figure 4 shows the relationship between the time (t) and the square of the distance from the LTO/LPS interface to the phase boundary (𝜉2) along the white arrows shown in Fig. 2c and Fig. 3c. Note that expanded view of the initial stage is shown in Fig. S9. A linear relationship was observed at each temperature for both Li insertion and extraction, and each value k was calculated from the latter part of the plots. This is because of the possibility that, at the beginning of the measurement, the phase boundary movement may be limited by the interfacial electrochemical reaction occurring at the LTO/LPS interface. If the interfacial electrochemical reaction determines the overall rate, the position of the phase boundary (𝜉2) would increase linearly with the square of the time (t2) and not with t. Table 1 shows the values of k at each temperature for both Li insertion and extraction. We analyzed 4 different particles (yielding 7 separate profiles) for Li extraction and 3 different particles (yielding 4 separate profiles) for Li insertion. From these data, we determined the parabolic rate constants of the phase boundary movement. As shown, the rate constant was 4.0 ± 0.7× 10−11 cm2/s at 30 °C and 1.9 ± 0.6 × 10−9 cm2/s at 105 °C for Li extraction and 3.6 ± 0.9 × 10−13 cm2/s at 30 °C and 3.2 ± 0.3 × 10−11 cm2/s at 105 °C for Li insertion, respectively. The value of k in graphite (from phase 2 to phase 1) at room temperature has been evaluated to be 3.23 × 10−11 cm2/s.32 Note that the stages number shows the number of empty graphene layers between two adjacent Li-filled layers. Therefore, the phase boundary movement in LTO during Li extraction is of the same order as the transition from stage 2 to stage 1 in graphite whereas the phase boundary movement in LTO during Li insertion is two orders of magnitude slower than the transition from stage 2 to stage 1 in graphite. Assuming that 𝛶 is approximately 0.5,29 the chemical Li diffusion coefficient is estimated to be roughly on the same order as k. This estimation is consistent with the chemical Li diffusion coefficient of 1.6 × 10−11 cm2/s in Li4Ti5O12 measured using chronoamperometry at room temperature.2Fig. 4. Relationship between the time (t) and the square of the distance from the LTO/LPS interface to the phase boundary (𝜉2) along the white arrows shown in (a) Fig. 2c (Li extraction) and (b) Fig. 3c (Li insertion). Temperature, T[°C] Parabolic rate constant of the phase boundary movement, k [cm2/s]  Li extraction Li insertion 105 1.9 ± 0.6 × 10-9 3.2 ± 0.3 × 10-11 90 7.8 ± 1.8 × 10-10 1.4 ± 0.3 × 10-11 75 4.1 ± 0.9 × 10-10 6.8 ± 0.7 × 10-12 60 2.0 ± 0.4 × 10-10 3.0 ± 0.5 × 10-12 45 1.0 ± 0.2 × 10-10 9.9 ± 1.3 × 10-13 30 4.0 ± 0.7 × 10-11 3.6 ± 0.9 × 10-13Table 1 Parabolic rate constant of the phase boundary movement (k) determined by operando STEM-EELS. The chemical Li diffusion coefficient (DⅡ) follows an Arrhenius-type temperature dependence, expressed by equation (3):where D0 is the pre-exponential factor, Ea is the activation energy for Li-ion diffusion in Li4Ti5O12 during Li extraction and Li7Ti5O12 during Li insertion, R is the gas constant, and T is the temperature. Using equations (2) and (3), equation (4) can be obtained:Fig. 5. Temperature dependence of the parabolic rate constant of the phase boundary movement (k). The error bars represent the standard deviation of the parabolic rate constants measured across the multiple particles. Ea is the activation energy of Li-ion diffusion. Figure 5 shows the temperature dependence of k during Li insertion and extraction. From the slope of the log k vs. 1/T plots, the activation energy for Li-ion diffusion was determined to be 0.49 eV for Li4Ti5O12 (associated with Li extraction) and 0.59 eV for Li7Ti5O12 (associated with Li insertion). The activation energies of Li-ion diffusion in Li4Ti5O12 obtained by density functional theory calculations and NMR are 0.30–0.48 and 0.45 eV, respectively,54,55 which are consistent with the value obtained in this study. The Arrhenius plot for the particle shown in Fig. 2 is presented in Fig. S10. The subscripts of each point in Fig. S10 show the order of the operando STEM-EELS measurements. Operando STEM-EELS experiments at 90 °C were conducted twice, once at the start and once at the end. The two values of k obtained at 90 °C were nearly identical (7.6 × 10−10 and 7.8 × 10−10 cm2/s), demonstrating the high reliability of our measurements. This result also confirms that electron beam damage does not significantly affect the kinetics of the phase boundary movement. Next, we compared the activation energy determined in this study with those of other processes. The activation energies for Li-ion diffusion in LiFePO4 and LiMn2O4 have been calculated to be 0.40 and 0.52 eV, respectively.25,56 Further, the activation energies of the charge-transfer resistance at the LiCoO2/lithium phosphorus oxynitride (LiPON) interface, LiCoO2/PC-based liquid electrolyte interface, and LTO/PC-based liquid electrolyte interface are 0.59, 0.64 and 0.50 eV, respectively.57–59 The relatively moderate activation energy of 0.49 eV for Li4Ti5O12 measured in this study suggests that the ionic diffusion and phase transition processes during Li extraction in LTO are not slow. Moreover, LTO has fast ionic-diffusion pathways along grain boundaries and defects.45 These factors contribute to its high-rate discharge performance, despite its two-phase reaction mechanism. However, Li7Ti5O12 has the higher activation energy of 0.59 eV than Li4Ti5O12. The result explains why Li insertion requires a significantly longer time than Li extraction. Our proposed temperature-resolved operando STEM-EELS technique can be used to investigate the ionic diffusion and phase boundary movement in these local regions, including surfaces, defects, and grain boundaries. We investigated the temperature dependent ionic diffusion using operando STEM-EELS. Our method enabled the real-space observation of the phase boundary movement in LTO at the nanometer scale, as well as the evaluation of the parabolic rate constant of the phase boundary movement at different temperatures. The rate constant was 4.0 ± 0.7 × 10−11 cm2/s at 30 °C to 1.9 ± 0.6 × 10−9 cm2/s at 105 °C for Li extraction and 3.6 ± 0.9 × 10−13 cm2/s at 30 °C and 3.2 ± 0.3 × 10−11 cm2/s at 105 °C for Li insertion, respectively. From the temperature dependence of the rate constant, the activation energy of Li-ion diffusion was determined to be 0.49 eV for Li4Ti5O12 (associated with Li extraction) and 0.59 eV for Li7Ti5O12​ (associated with Li insertion). This relatively low activation energy in Li4Ti5O12 is one reason why LTO achieves high-rate discharge performance, despite its two-phase character.EXPERIMENTAL METHODSFabrication of the Solid-State Li-Ion Battery. Commercially available LTO particles with an average particle size (D50) of approximately 4 μm (Toshima Co., Ltd.), LPS particles with a D50 of approximately 1 μm (PO0119, MSE Supplies LLC), and In-Li foil with a thickness of 150 µm were used as the working electrode, solid electrolyte, and counter electrode, respectively. The ionic conductivity of LPS was approximately 1.0 × 10−3 S/cm at room temperature. The thickness of the solid electrolyte layer was approximately 400 µm. The cell was pressed at 300 MPa and 120 °C in an argon-filled glove box with a dew point below −70 °C.Electrochemical Analyses. The charge and discharge measurements shown in Fig. 1c were conducted in a transmission electron microscope at a constant current of 300 nA, followed by constant voltages of 0.4 and 1.6 V vs. In-Li using a potentio/galvanostat (SP-200, BioLogic Science Instruments). The charge and discharge measurements were terminated when the currents decreased to approximately 10 nA. The EIS spectra shown in Fig. 1d were also obtained in a transmission electron microscope, covering a frequency range from 3 MHz to 100 mHz. The chronoamperograms shown in Fig. 1e were obtained at 1.6 and 0.4 V vs. In-Li, and the currents were recorded at intervals of 0.05 s. The battery sample was heated using a ceramic heater built into the TEM holder. The temperature was measured using a thermocouple and maintained at a constant temperature through feedback control. After the thermocouple temperature stabilized, the EIS measurements were repeated. Once the semicircular arcs in the high-frequency region, representing the resistance of LPS, became constant, chronoamperometry and operando STEM-EELS measurements were conducted. The temperature was controlled with a precision of ± 0.005 °C.Sample Preparation for STEM-EELS. The battery cell was placed on a vacuum-transfer FIB holder inside the glove box and then transferred to a FIB system (FB-2100, Hitachi High-Tech. Corp.) without exposure to air. The LTO side of the battery was thinned using a FIB accelerated at 40 kV, resulting in a region with a thickness of 100–200 nm.STEM-EELS. A 300-kV transmission electron microscope (JEM-ARM300F2 GRAND ARM2, JEOL Ltd.) equipped with an EEL spectrometer (Gatan imaging filter Continuum-K3, Gatan Inc.) was used for operando STEM-EELS observations. The accelerating voltage was 300 kV, electron probe current was 71–529 pA. Operando STEM-EELS data were acquired using the in-situ spectral imaging (SI) mode of GMS 3.60 in Digital Micrograph (Gatan Inc.), enabling sequential STEM-EELS measurements without dead time between the STEM-EELS SIs. Detailed information on the design of the vacuum-transfer TEM holder with biasing electrodes is available in the literature.60ASSOCIATED CONTENTSupporting InformationThe Supporting Information is available free of charge on the ACS Publications website.Movie of the Li-ion diffusion at different temperatures (MP4), Electron diffraction (Fig. S1), Core-loss EELS analysis of the Ti-L and O-K edges (Fig. S2), Reproducibility and representativeness for Li extraction (Fig. S3, S4), Comparison between diffusion-limited and reaction-limited regime (Fig. S5), Formation mechanism of core-shell structure during Li insertion (Fig. S6), Reproducibility and representativeness for Li insertion (Fig. S7, S8), Relationship between the time (t) and the square of the distance from the LTO/LPS interface to the phase boundary (𝜉2) (Fig. S9), Reliability of operando STEM-EELS and effect of electron irradiation damage (Fig. S10).AUTHOR INFORMATIONCorresponding AuthorYuki Nomura − Nanostructures Research Laboratory, Japan Fine Ceramics Center, 2–4–1 Mutsuno, Atsuta-ku, Nagoya, Aichi, 456–8587, Japan; orcid.org/0000- 0002-6091-5902; Email: y_nomura@jfcc.or.jpAuthorsKazuo Yamamoto − Nanostructures Research Laboratory, Japan Fine Ceramics Center, 2–4–1 Mutsuno, Atsuta-ku, Nagoya, Aichi 456–8587, Japan; orcid.org/0000-0001-7651-943XNaoaki Kuwata − Research Center for Energy and Environmental Materials, National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan; orcid.org/0000-0002-0736-6967Tsukasa Hirayama − Nanostructures Research Laboratory, Japan Fine Ceramics Center, 2–4–1 Mutsuno, Atsuta-ku, Nagoya, Aichi 456–8587, JapanNotesThe authors declare no competing financial interest.ACKNOWLEDGMENTWe thank Ms. Misaki Hasegawa for the support of sample preparation for STEM-EELS observation. This work was partly supported by the Innovative Science and Technology Initiative for Security (JPJ004596) from the Acquisition, Technology & Logistics Agency, the Grant-in-Aid for Scientific Research KAKENHI (23K13837, 23H00241, 23H01858) from the Japan Society for the Promotion of Science, the GteX Program (JPMJGX23S2) from the Japan Science and Technology Agency, and the Kazato research Foundation.REFERENCES(1) Ohzuku, T.; Ueda, A.; Yamamoto, N. Zero-Strain Insertion Material of Li[Li1/3Ti5/3]O4 for Rechargeable Lithium Cells. J. Electrochem. Soc. 1995, 142 (5), 1431–1435.(2) Takami, N.; Hoshina, K.; Inagaki, H. Lithium Diffusion in Li4/3Ti5/3O4 Particles during Insertion and Extraction. J. Electrochem. Soc. 2011, 158 (6), A725–A730.(3) Zhu, G.-N.; Wang, Y.-G.; Xia, Y.-Y. Ti-Based Compounds as Anode Materials for Li-Ion Batteries. Energy Environ. Sci. 2012, 5 (5), 6652–6667.(4) Yi, T.-F.; Yang, S.-Y.; Xie, Y. 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