1. Introduction
Lithium-ion (Li-ion) batteries are crucial in the current digital electronic era, powering a vast range of devices, from handheld electronics to electric vehicles and large-scale energy storage systems [1–3]. Despite their widespread adoption, Li-ion batteries have significant safety risks, notably thermal runaway and ignition [4]. To address these concerns, researchers have focused on the development of safer alternatives, such as batteries with aqueous electrolyte systems. However, these alternatives typically have lower energy densities and operational voltages because of their limited electrochemical windows [5].
Recently, Suo et al. introduced a “water-in-salt” (WIS) electrolyte that substantially widens the electrochemical window of aqueous batteries, enabling higher voltage operations while enhancing safety by reducing flammability [6,7]. Nevertheless, this breakthrough led to challenges in controlling electrolyte reactivity to prevent harmful side reactions, especially at the electrode interface. To resolve this issue, Wang et al. proposed a hybrid electrolyte with aqueous and non-aqueous components that generates a solid-electrolyte interface (SEI) on the anode [8], which is crucial for improving Li-ion battery safety and efficiency by preventing electrode degradation and electrolyte decomposition. The SEI comprises organic and inorganic phases, and LiF is a particularly effective phase for passivating electrodes and boosting cycling stability [1,3,9]. The microstructure and composition of the SEI have key influences on battery performance. For example, a porous SEI can lead to continuous breakdown and repair cycles [10], inducing side reactions and reducing battery life. Conversely, a dense, highly crystalline SEI remains intact during charge–discharge cycles, minimizing side reactions and enhancing cyclic performance [11].
Even though the mechanism of SEI formation, its composition, and its influence on battery performance have been extensively studied [1,11], all the SEI characteristics are not fully understood. Therefore, sophisticated analytical methods are required to examine the complex structure and function of SEI layers.
In this study, we used high-resolution transmission electron microscopy (HRTEM) to examine the microstructure of hybrid aqueous/non-aqueous electrolytes and their role in stabilizing SEI formation on lithium titanate (Li4Ti5O12; LTO). We also performed atomic-scale analysis of the interactions of the LiF SEI layer and its stability on the LTO electrode through density functional theory (DFT) calculations. We believe that this study contributes to the development of safer and more reliable batteries.
2. Experimental
2.1 Electrode and Electrolyte preparation
Cathode electrodes comprising LiMn2O4 (LMO, L&F Co., Ltd.) were produced by blending the active material, carbon black (Super C, Timcal), and polyvinylidene fluoride (PVDF, Solvay) in a 90:5:5 weight ratio with N-methylpyrrolidinone (NMP, 99.5%, Sigma-Aldrich). The mixture was spread onto 20-μm-thick titanium foil and dried in vacuum at 100°C for 12 h. The same procedure was followed to prepare anode electrodes comprising Li4Ti5O12 (LTO, POSCO Future M), except for the use of aluminum foil as the substrate. The hybrid aqueous electrolyte (HT-29) was formulated by combining lithium bis(trifluoromethanesulfonyl)imide (LiTFSI, 99.95%, Sigma-Aldrich), water, and trimethylphosphate (TMP, 99%, Sigma-Aldrich) in a 37:40:23 molar ratio [12].
2.2 Electrolyte characterization
The ionic conductivity of the electrolyte was determined by a conductivity meter (S213-meter, Mettler-Toledo), and the viscosity of the electrolyte was determined using a rotational viscometer (Brookfield DV3TLV). The linear sweep voltammetry (LSV) tests for the electrolyte were performed with a three-electrode system constituting an inactive current collector (titanium foil), activated carbon, and Ag/AgCl as the working, counter, and reference electrodes, respectively. The LSV measurements were conducted on a Solatron 1260 instrument equipped with a 1287 Electrochemical Interface (Solatron Metrology, UK) at a scanning rate of 0.1 mV s−1.
2.3 Cell assembly and Testing
The cathode, anode, aqueous electrolyte, and separator were assembled in a dry room with a controlled relative humidity of <0.2%, as depicted in Fig. 1. A 260-μm-thick Whatman GF/A glass fiber was used as the separator. A standard 2032-coin cell was filled with 100 μL of electrolyte, sealed, and allowed to stand for at least 12 h to ensure thorough impregnation. Charge–discharge testing was conducted in a temperature-controlled environment using a Maccor 4000 battery tester, cycling between 2.8 and 1.5 V at a 0.5C rate (0.5 mA cm−2).
2.4 Microstructural analysis
The post-cycling crystallographic orientation of the SEI and LTO was determined using field-emission transmission electron microscopy (FE-TEM, JEOL, JEM-2100F) at 200 kV. Samples were prepared with a focused ion beam (FIB, JEOL, JIB-4601F). HRTEM images and fast Fourier-transform (FFT) analysis were used to verify the crystal structure of the SEI on the LTO electrode. The chemical composition of the SEI was further analyzed using X-ray photoelectron spectroscopy (XPS, Axis-Supra, Kratos).
2.5 Computational details
Spin-polarized DFT calculations were conducted using the Vienna Ab initio Simulation Package (VASP) [13] using the projector-augmented wave (PAW) method and the Perdew–Burke–Ernzerhof (PBE) [14] exchange-correlation functional within the generalized gradient approximation (GGA). The Grimme D3 dispersion correction was employed to account for van der Waals interactions [15], and calculations were performed with a plane-wave basis set with a cutoff energy of 600 eV for increased accuracy.
For bulk structure optimization, Monkhorst–Pack (6×6×6) k-point meshes for LiF and (3×3×3) for LTO were used. Slab models representing the (200) crystallographic planes of LiF and LTO were derived from these optimized bulk structures, selected based on the results of TEM analysis. The slabs were further optimized with (6×6×1) k-point meshes for LiF and (3×3×1) for LTO. The lowest-energy configurations from various (200) slab models with different surface terminations were chosen for subsequent analyses, as shown in Fig. 2 and 3. The LiF (200)/LTO (200) interface model, constructed from these slabs, was refined using a (6×3×1) k-point meshes, as illustrated in Fig. 4. A 20 Å vacuum layer was introduced perpendicular to the slabs to prevent artificial periodic interactions, and the bottom two layers were fixed to mimic the bulk conditions, allowing the upper layers to relax.
The chemical stability and chemical bonding at the LiF/LTO interface were assessed by calculating the formation energy (Ef) as follows:
where ELiF/LTO, ELTO, and ELiF represent the total energies of the combined interface, LTO layer, and LiF layer, respectively.
3. Results and Discussion
3.1 Cell performance
Table 1 shows the ion conductivity of the hybrid aqueous electrolyte at 23C and its viscosity at 25C and 40C. The ion conductivity of this hybrid aqueous electrolyte was smaller than that of a non-aqueous electrolyte [16]. Its viscosity was higher than that of a traditional non-aqueous electrolyte, owing to the high concentration of LiTFSI salts in the former.
The electrochemical stability window (ESW) of the hybrid aqueous electrolyte was determined from the linear sweep voltammetry curves (Fig. 5). The cathodic and anodic limits were 1.76 and 4.66 V, respectively, with an overall ESW of 2.9 V. The redox reaction potentials of LMO and LTO are completely inside the ESW of the hybrid aqueous electrolyte.
An LMO–LTO full cell was constructed to assess the electrochemical performance of the hybrid aqueous electrolyte. The initial voltage profile, recorded at a current density of 0.5C (0.5 mA cm−2), is shown in Fig. 6. The first cycle exhibited a charge voltage plateau at ~2.50 V and a discharge plateau at 2.35 V, indicative of Li+-ion insertion and removal in the LMO–LTO system. The initial discharge capacity (146.1 mAh g−1) was marginally lower than the theoretical capacity (175 mAh g−1). The coulombic efficiency of the first cycle was 80.1%, suggesting electrolyte consumption for SEI formation and H2O decomposition at the anode. In the subsequent cycle, the coulombic efficiency improved to above 90%, suggesting minimal interface formation. After 100 cycles at 0.5C, the cell retained a discharge capacity of 51.9 mAh g−1, which is 35.4% of the initial capacity, and a coulombic efficiency of ~97.8%.
The XPS spectra of the SEI on the anode are shown in Fig. 7. The F 1s spectrum of the LTO surface shows peaks of −CF2 (686.3 eV) of the PVDF binder in the electrode, −CF3 (688.6 eV) of the TFSI− salt anion, and the inorganic LiF (685.3 eV) formed by the reduction of [Li]2+[TFSI]−. This suggests that the LiF molecules within the SEI effectively inhibit the hydrogen evolution reaction on the anode by creating a protective film on the LTO surface.
Additionally, the C 1s spectrum exhibited peaks of −CF3 (293 eV) from TFSI− or its reduction byproducts, −CO3 (290.5 eV) originated from the reaction of CO2 with H2O, alkyl carbonates (289 eV), ether (C–O) (286 eV), and other carbonaceous sp3- or sp2-hybridized elemental carbon species below 284 eV. The −CO3 peak is indicative of the interaction between atmospheric CO2 and H2O [17].
3.2 Microstructure of the SEI
Fig. 8 shows a high-resolution image of the specimen after 30 cycles, in which the SEI layer is discernible by the fringe contrast with the LTO electrode, measuring ~10 nm in thickness, as shown in Fig. 8a. Before FIB sample preparation, a Pt layer was applied to the SEI surface to prevent ion-beam damage [18]. The FFT analysis (Fig. 8b) identified the crystal structure of the SEI, as LiF has a face-centered cubic (FCC) lattice (Fm-3m; a = 4.027 nm). The epitaxial growth of the SEI on the LTO electrode is facilitated by the minimal lattice mismatch and the similarity in the FCC crystal structures of the LTO and LiF phases. This allows the LTO to serve as a substrate for the LiF phase. The results confirm that the LiF SEI layer grew epitaxially on the LTO electrode surface, for which the (200) and (111) planes are parallel between the two phases. HR-TEM observations revealed a semi-coherent interface between the LTO and LiF SEI layers, as indicated by the nearly 2:1 ratio of their lattice constants. This small lattice mismatch suggests that the LTO electrode is an effective nucleation site for the LiF SEI layer. The well-formed SEI layers prevent continuous side reactions and repeated SEI formation, leading to reduced resistance and enhanced stability [11]. The crystalline and dense SEI layer fully passivates the electrode surface, maintaining mechanical stability during cycling. The epitaxially grown LiF SEI on the LTO electrode demonstrates uniformity and crystallinity with minimal defects, effectively isolating the electrode from the reactive aqueous electrolyte and providing a robust surface.
3.3 DFT calculations
DFT calculations were conducted to investigate the atomic-scale interactions and stability of the LiF SEI layer on the LTO electrode. The results indicate an energetically favorable LiF-LTO interface structure with a formation energy of −5.26 eV, suggesting strong adhesion and a thermodynamically stable interface. Specifically, upon the formation of the interface, the bond length of Li–F within the LiF(200) surface decreased from 2.04 to 1.97 Å. Additionally, the bond lengths of Ti–O and Li–O within the LTO (200) surface changed from 1.93 to 2.09 Å and 1.97 to 1.92 Å, respectively. Moreover, as a direct consequence of the interface formation, new Li–O and Ti–F bonds with bond lengths of 1.98 and 2.13 Å, respectively, have emerged (Fig. 4). Furthermore, the lattice mismatch between LiF and LTO is only 2.22%, suggesting the epitaxial growth of the SEI layer on the LTO substrate. Significant charge redistribution was observed at the LiF/LTO interface, which is critical for the stability of the SEI (Fig. 9). The charge density difference (Δρ) for LiF/LTO was calculated as follows:
where ρ(LiF/LTO), ρLTO, and ρLiF represent the charge densities of the interface, the LTO layer, and the LiF layer, respectively. The results of these calculations indicate electron redistribution and localization because of interfacial interactions. In addition, Bader charge analysis [20] revealed that the LTO layers gain 0.30 electrons, whereas the LiF layers lose 0.30 electrons, which enhanced the ionic character of the bonding and promoted stability at the LiF/LTO interface.
4. Conclusions
In this study, a LTO/LMO full cell was fabricated with a hybrid aqueous electrolyte. The cell showed superior discharge performance during 0.5C cycle tests. TEM revealed the epitaxial growth of the LiF SEI layer on the LTO electrode, in which the (200) and (111) planes were aligned between the two phases. This structural conformance results in lattice coherence at the interface, which significantly improves the electrochemical stability and functionality of the cell. HRTEM revealed a semi-coherent interface, suggesting that the LTO substrate served as a nucleation site for the LiF SEI layer. DFT calculations further confirmed the stability and favorable energetics of the formation of an LiF SEI layer on the LTO substrate which is consistent with the experimental evidence of epitaxial growth. The synergistic use of experimental and computational methods in this study contributes to the development of LTO/LMO full cells, especially for enhancing both discharge capacity and stability. The DFT calculations and HRTEM findings provide an in-depth understanding of the growth mechanisms of the SEI layer and its influence on battery performance. This study highlights the necessity of an integrated approach to refining lithium-ion batteries for the future.
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