Investigation of Zr doping and Electrochemical study on Li1+xTa1–xZrxSiO5 (x = 0–0.3) Solid Electrolyte for Lithium-Ion Battery
Article information
Abstract
Materials with high lithium-ion conductivity are promising for use as electrolytes in solid-state batteries. However, several technical challenges, such as stability in air and optimal synthetic conditions, inhibit their extensive application. In this study, LiTaZrSiO5 solid electrolytes (SEs) are prepared at different sintering temperatures (900–1100°C) and characterized using X-ray diffraction, scanning electron microscopy, and electrochemical impedance spectrometry. Correlations among the sintering temperature, structural properties, and ionic conductivity of the LTSO SEs are examined systematically and discussed. According to the results, high ionic conductivity can be achieved by optimizing the sintering temperature of the SEs. The LTSO SE sintered at 1050°C exhibited the highest ionic conductivity. To further improve ionic conductivity, the sintering time was optimized in the range of 6–24 h. The conductivity improved as the sintering time increased from 6 to 12 h, and the best conductivity was achieved at 12 h. For sintering durations longer than 12 h, the conductivity of the material decreased. Further, doping emerged as a prominent strategy for increasing the ionic conductivity of the SEs. The effects of different concentrations of Zr dopant on the physio-chemical (structural, morphological) and electrochemical properties of the Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs were investigated. Compared to the results obtained for the other SEs, extraordinary results were obtained by doping the Li1.1Ta0.9Zr0.1SiO5 SE with zirconium, such as increased ionic conductivity, reduced charge transfer resistance, reduced crack formation, and increased SE particle size. In summary, the sintering temperature, sintering time, and dopant concentration in the synthesized Li1+xTa1–xZrxSiO5 (x = 0–0.3) precursors should be optimized to promote the development of LTZSO electrolytes suitable for use in future SSBs.
INTRODUCTION
Development of advanced sustainable lithium-ion batteries (LIBs) is necessary to mitigate the problems of environmental degradation and greenhouse effect, and reduce the consumption of natural energy resources [1–5]. In recent decades, LIBs have been extensively used in portable electronics, electric vehicles, electronic gadgets, and energy storage devices because of their high energy density, storage capacity, long lifespan, and light weight [1–4,6–12]. Nevertheless, typical LIBs continue to be affected by safety hazards, such as Li-dendrite formation, explosions resulting from the use of combustible organic solvents, electrolyte leakage, and undesirable side reactions [8,9]. Therefore, the use of solid electrolytes (SE) has been investigated to solve the safety issues caused by the leakage of flammable organic liquid electrolytes (LEs) in conventional LIBs [1,7,9,13–15]. Solid-state batteries (SSBs) have attracted considerable interest in recent years, especially for use in modern electric vehicles (EVs). Among the different materials used in such batteries components, the SE is the most reactive and complex material, and it considerably influences SSB failure. Therefore, the development of a highly conductive and stable SE is critical to the success of SSBs.
Many studies have focused on improving the structural stability and conductivity performance of SEs. Generally, garnet Li7La3Zr2O12, NASICON (Li1+xAlxTi2–x(PO4)3, LATP), perovskite-type (Li3xLa0.67–xTiO3, LLTO), Li10GeP2S12 family, and argyrodite sulfide-based SEs are used in SSBs [16–23]. However, many of these SEs are unstable against air or Li metal [1,7,13,24]. Therefore, new Li-ionic conductors that offer both remarkable ionic conductivity and excellent stability need to be developed as SEs for SSBs. To achieve high ionic conductivity, from the crystal structure perspective, research is being conducted to achieve the lowest Li-ion migration barrier [25–27]. For example, materials with a body-centered cubic (BCC)-like anion framework structure exhibit high ionic conductivity owing to the low activation energy required for Li-ion movement [25]. However, despite considerable research, the only materials with a BCC-like anion framework structure are LGPS and Li7P3S11, which are sulfide-based SEs [25]. Although, garnet- and NASICON-structure-based oxide SEs are grouped under non-BCC-like anion framework materials, they provide higher ionic conductivity [25,26]. This is because certain element doping compositions within the crystal can induce a disordered Li-ion sublattice that helps ions move quickly by using the repulsive forces between adjacent ions [26,27]. However, the characteristic of the crystal structure framework of SEs that causes this higher ionic conductivity remains controversial.
Recently, Mo et al. reported a computational model of fast ion diffusion in Li superionic conductors (LICs) to explain the concerted migration of multiple Li ions [26]. Theoretically, this model is based on the strategy of inserting Li-ions into high-energy sites to initiate concerted ion migration with a lower energy barrier [26]. Based on this knowledge, Chen et al. and Guo et al. prepared a new SE called LiTaSiO5 (LTSO) by using quenching and the solid-state method [24,28]. Chen et al. studied the structure and Li-ion diffusion mechanism of Li1+xTa1–xZrxSiO5 (0.5 ≥ x ≥ 0) by using the ab-initio molecular dynamic (AIMD) simulation method [24]. Guo et al. studied the crystal structure, electrical properties, and Li-ion migration pathways of LiTaSiO5 and Li1.1Ta0.9Zr0.1SiO5 by performing AIMD stimulations [28]. They proposed a monoclinic structure with the space group P 21/c (No.14) [29]. According to the literature, the LiTaSiO5 crystal structure is composed of SiO4 tetrahedron and a corner-sharing TaO6 octahedron. The Li-ions occupying the 4e tetrahedral sites and the LiO4 tetrahedrons share edges with adjacent SiO4 and TaO6 polyhedrons [24,26,28,29]. This 3D framework with a suitable bottleneck size is suitable for Li-ion transfer from the perspective of accomplishing concerted Li-ions migration. Furthermore, these studies demonstrated that the high ionic conductivity of the material at room temperature is due to zirconium (Zr) doping. The Zr4+ (0.72 Å) ionic radius is higher than Ta5+ (0.64 Å) and it assists to increase the lattice parameters of the SEs [24,30]. Additionally, the excess lithium is added to composition to compensate the doping of Zr4+ instead of Ta5+. The excess lithium is beneficial for the ionic conductivity of the electrolyte.
The preparation methods used in these works were almost similar to the solid-state reaction described in previous report [28]. Nonetheless, comprehensive and extensive focus on the preparation of LTSO SE is highly desirable for comprehending the effect on the phase formation and ionic conductivity of LTSO. In the present work, we report the preparation of Li1+xTa1–xZrxSiO5 (x = 0–0.3) by using the conventional solid-state method. The prepared LTSO SE is sintered at different temperatures and for different durations to evaluate the effects of these variables on its structure, morphology, and ionic conductivity. We systematically investigate the effect of sintering temperature, time, and doping concentration on the structural formation, surface morphology and conductivity performance of the prepared LTSO SEs.
EXPERIMENTAL
Preparation of Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs
Li1+xTa1–xZrxSiO5 (x = 0–0.3) compounds were prepared using a ball-milling method (PULVERISETTE 6 Planetary Mono Mill). In this process, lithium carbonate (Li2CO3, Sigma Aldrich, USA), tantalum pentoxide (Ta2O5, DAEJUNG, Korea), zirconium oxide (ZrO2, Sigma Aldrich, USA), and silicon oxide (SiO2, Sigma Aldrich, USA) were used as the starting materials. To compensate for lithium loss at high temperatures during synthesis, the compounds were prepared with 5 wt.% excess Li2CO3. The precursors were mixed in stoichiometric amounts with ethanol in a ball milling jar and wet mixed at 200 rpm for 6 h until a fine homogeneous mixture was obtained. Thereafter, the solvent was evaporated, and the mixture was dried in an oven at 50°C for 12 h. The powders obtained after drying were calcined at 650°C for 12 h in air atmosphere. The calcined powders were cold pressed under a pressure of 100 MPa to obtain pellets with a thickness of approximately 1 mm and diameter of 10 mm. Lastly, the final product Li1+xTa1–xZrxSiO5 (x = 0–0.3) was obtained by sintering the pellets at 900°C for 6–24 h in air atmosphere.
Material and electrochemical characterizations of Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs
A high-resolution X-ray diffraction (HR-XRD) analyzer was used to determine the crystal structure of the sintered sample. The Bragg’s equation can be used to calculate the lattice spacing (d) of the electrolyte [31].
Where the λ is the X-ray wavelength (0.15418 nm) and θ is the Bragg’s diffraction angle. The average grain size of the electrolyte is calculated using the Debye-Scherrer equation [4,31].
K is the Scherrer constant (0.9) and β is the full-width at half maximum (FWHM).
The surface morphology and particle size of the synthesized Li1+xTa1-xZrxSiO5 (x = 0–0.3) SE were confirmed using field emission scanning electron microscopy (FE-SEM). Inductively coupled plasma luminescence analysis (ICP-OES, Inductively Coupled Plasma Optical Emission Spectroscopy) was performed to study the stoichiometry of the sample. Its ionic conductivity was measured using an electrochemical impedance spectrometry analyzer (EIS, LCR meter 4284A, HP Co., USA) in the frequency range of 1 MHz–25 Hz. A thin layer of Pt was sputtered on each surface of the sintered pellets as an electrolyte. Ion conductivity was calculated using the equation σ = d/(A × R), where d is the pellet thickness (0.1 cm), A is the contact area of the pellet, and R is the total resistance value measured using EIS.
RESULTS AND DISCUSSION
Optimization of sintering temperature of LiTaSiO5 (LTSO) SE
Analysis of structural characteristics of LiTaSiO5 (LTSO) SE by optimizing the sintering temperature
Sintering conditions such as temperature, time, and atmosphere of the sintering process are important factors affecting the structural and electrochemical characteristics of SEs [32]. Therefore, it is important to optimize the sintering conditions for synthesizing the SE. In addition, such optimization is necessary to prevent unwanted side effects, such as Li-loss during the sintering process and optimization of sample composition. As the primary heat treatment, LiTaSiO5 SE pellets were sintered in a box furnace in air atmosphere for 6 h at temperatures of 900°C to 1100°C, where the temperature was increased in steps of 50°C, to check the optimum synthesis conditions.
X-ray diffraction analysis was performed to determine the structural characteristics of the samples sintered under various sintering temperature conditions. Fig. 1a depicts the XRD patterns of the LTSO SEs sintered at different temperatures ranging from 900°C to 1100°C. The XRD patterns of the sample sintered at 900°C and 950°C did not match the reference pattern (ICSD#39648) [27]. This result was attributed to the failure to bind the oxygen atoms and silicon atoms during synthesis of the LiTaSiO5 SE under low-temperature conditions. As depicted in Fig. 1a, the diffraction peaks of the pristine LTSO powder at 1000°C or higher belong to a monoclinic phase, confirming that the LTSO structure clearly changes from the trigonal phase (space group: P 1, LiTaO3) to the monoclinic phase with increasing sintering temperature. This result indicated that the sintering temperature had a significant effect on the crystal structure and phase formation of the LTSO SE. Additionally, a minor impurity phase (LiTaO3) was detected in all LTSO powders sintered at high temperatures.
(a) XRD patterns of LiTaSiO5 SEs at different synthetic temperatures, (b) graph of the (110) plane intensity, (c) d (110) plane spacing of LiTaSiO5 SEs at different temperatures, and (d) EIS curves of LiTaSiO5 SEs prepared/sintered at different temperatures.
Fig. 1b depicts the XRD pattern of the (110) plane of the LTSO SE obtained by sintering at 1000–1100°C. For the pellets sintered at 1000–1050°C, the (110) peak intensities increased. However, for the pellets sintered at 1100°C, the peak intensities and d-spacing values decreased (Fig. 1c). This indicated that LTSO might have undergone lithium evaporation at around 1100°C. For clarity, the lattice parameters of the samples sintered at different temperatures are summarized in Table 1. The lattice parameters and volume of the LTSO SE were higher only at the sintering temperature of 1050°C. Seemingly, the lattice parameters of the sample sintered at 1100°C were marginally lower, possibly because of lithium loss.
Morphological and electrochemical characteristics of LiTaSiO5 (LTSO) SE by optimizing the sintering temperature
Electrochemical impedance spectroscopy (EIS) was performed to study the ionic conductivity of the LiTaSiO5 SE as a function of the sintering temperature in the range of 900–1100°C. The Cole–Cole plots of all the sintered samples exhibit only one semicircle in the high-to-low-frequency region (Fig. 1d). The semicircle in the graph is related to the total impedance of grains and grain boundaries. Compared to the resistance of the low-temperature sample, the resistance of the high-temperature samples was lower, and it decreased and then increased as the sintering temperature increased. Li evaporation during long-term sintering may have led to an increase in charge transfer resistance [4,33]. The ionic conductivity of the SE samples at RT increased as the sintering temperature increased, and it reached approximately 7.76×10–8 S cm–1 for the LTSO SE sintered at 1050°C. However, the ionic conductivity of the sample sintered at 1100°C was lower. The sample sintered at 1050°C exhibited the smallest grain boundary resistance value (2,090,000 Ω) in the EIS curve, and this result was attributed to its higher d-spacing value and lattice parameter, as confirmed by XRD analysis (Fig. 1a–c). Hence, the sintering temperature of 1050°C was selected as a suitable condition for further synthesis of LTSO SEs.
FE-SEM was used to examine the surface morphology of the LTSO SEs pellets. Fig. 2a–d present morphological images of the LTSO pellets prepared by sintering at 900–1100°C. The images in Fig. 2a,b confirm that sintering at 900–950°C was inadequate for particle binding. According to Fig. 2c, at the sintering temperature of 1000°C, the interparticle contact area increased due to neck growth, and a continuous pore channel was developed in the pellets [34–37]. At the sintering temperatures of 1050–1100°C, this continuous pore channel vanished owing to particle fusion, and large pores in the pellets were eliminated because the particle size increased (Fig. 2d) [34–37]. Additionally, cracks were observed in all the pellets sintered at high temperatures. Hence, sintering at temperatures exceeding 1000°C was confirmed to have increased the pellet’s compaction strength and integrity.
FE-SEM images of LiTaSiO5 SEs prepared by sintering at (a) 900°C, (b) 950°C, (c) 1000°C, (d) 1050°C, and (e) 1100°C.
Optimization of sintering time of LiTaSiO5 (LTSO) SE
Analysis of structural characteristics of LiTaSiO5 (LTSO) SE by optimizing the sintering time
The XRD patterns of the LTSO SEs sintered at 1050°C for durations of 6 h to 24 h are presented in Fig. 3. These XRD results exhibit characteristic peaks, and there are no significant differences in the crystallinity of the samples prepared by sintering for 6 h to 18 h (Fig. 3a). Table 2 summarizes the lattice parameters of the samples prepared by sintering for different durations. As the sintering time increased from 6 h to 12 h, the average grain size of the samples increased owing to grain boundary migration [32,38]. In samples sintered for more than 12 h, the average grain size decreased owing to enhanced grain boundary generation [39]. Table 3 summarizes the average grain sizes of the samples sintered for different durations. The peak intensity of the sample sintered for 24 h decreased owing to Li evaporation (Fig. 3b). In addition, LiTaO3 impurity peaks were identified in the spectra of all the prepared samples owing to inadequate bonding between oxygen and silicon atoms.
(a) XRD patterns of LiTaSiO5 SEs sintered for different durations, (b) graph of (110) plane, (c) d (110) plane spacing of LiTaSiO5 SEs sintered for different durations, and (d) EIS curves of LiTaSiO5 SEs sintered for different durations.
Morphological and electrochemical characteristics of LiTaSiO5 (LTSO) SE by optimizing the sintering temperature
Fig. 3d shows the Cole–Cole plot of the LTSO SEs sintered for different durations (6–24 h). In Fig. 3d, a correlation is established between the sintering time and resistance properties of the LTSO SEs. As the temperatures increased, the diameter of the semicircle decreased steadily, and simultaneously, ionic conductivities of the LTSO SEs increased. The total resistance (Rtotal) of the LTSO SEs decreased as the sintering time increased from 6 to 12 h. The total resistance of the LTSO SEs first decreased from 6 h to 12 h then increased as the holding time reached 18 h, as can be inferred from their microstructural properties and grain size data (Table 3). The samples sintered at 1050°C for 12 h exhibited the highest ionic conductivity (1.78×10–7 S cm–1), which was attributed to its lowest total resistance (1,350,000 Ω), highest d-spacing (110 plane) value, and improved crystallinity (Fig. 3b,c). Apart from sintering temperature, sintering time, too, significantly influenced Li+ volatilization. As the sintering time increased (> 18 h), the volatilization of Li+ in the LTSO SEs was unavoidable, which could have led to a decrease in peak (110) intensity (Fig. 3b). When the holding time was short (12 h), the conductivity of the LTSO SEs increased owing to the large average grain size and low intensity of the LiTaO3 phase. However, the conductivity of the LTSO SEs decreased as the holding time increased beyond a certain threshold (>12 h), probably because of Li+ volatilization and an increase in LiTaO3 phase intensity (Fig. 3a). Accordingly, prolonged sintering times adversely affected the ionic conductivity of the SEs. Therefore, the sintering time of 12 h was selected as the optimal condition for further processing.
The morphological images presented in Fig. 4 systematically reveal the effect of sintering time on the sintered LTSO SE pellets. Additionally, the porosity of the LTSO pellets decreased significantly as the sintering holding time increased. No major transformation in particle shape occurred with increasing sintering time. Moreover, all the samples exhibited crack formation owing to thermal expansion of particles.
FE-SEM images of LiTaSiO5 SEs sintered for (a) 6 h, (b) 12 h, (c) 18 h, and (d) 24 h.
Optimization of doping concentration of Li1+xTa1–xZrxSiO5 (x = 0–0.3) SE
Analysis of structural and morphological characteristics of Li1+xTa1–xZrxSiO5 (x = 0–0.3) SE by optimizing the doping concentration
An ICP-OES analysis was performed to check the actual stoichiometry of the Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs, and the results are summarized in Table 5. The ICP-OES analysis was carried out after pre-treating the SE samples with aqua regia. For five SE compositions, most of the target values of the synthesis materials and the values obtained through analysis were consistent.
ICP-OES analysis results of Li1+xTa1–xZrxSiO5 (x = 0–0.3) materials
The XRD patterns of the Zr-doped Li1+xTa1–xZrxSiO5 (x = 0–0.3) powders are presented in Fig. 5a. In the composition range of x = 0–0.2, the diffraction patterns were indexed based on the reference pattern (ICSD#39648) with a small amount of LiTaO3 impurity phase. For x = 0.1 or higher, the major (110) diffraction plane shifted to a lower degree, meaning that Zr4+ was successfully incorporated into the crystal structure, and the d-spacing and lattice parameter values increased, which favored fast Li-ion transfer (Fig. 5b,c). Table 4 summarizes the lattice parameters of the prepared Li1+xTa1–xZrxSiO5 (x = 0–0.3) electrolyte. The diffraction peaks shifted to lower angles owing to the larger ionic radius of Zr4+ (0.72 Å) than that of Ta5+ (0.64 Å) [24,30]. However, as the zirconium substitution composition value x increased, the LiTaO3 peak intensified, which was attributed to destabilization of the binding of oxygen and silicon due to zirconium substitution. Additionally, for x = 0.3, a new ZrSiO4 peak evolved in the sample [40,41]. Therefore, the zirconium substitution composition range for maintaining the LiTaSiO5 crystal phase was set to x = 0–0.2. The morphology of the as-synthesized Zr-doped Li1+xTa1–xZrxSiO5 (x = 0–0.3) SE powders sintered at 1050°C for 12 h in the air is depicted in Fig. 6. zirconium content played a critical role in decreasing cracking and porosity of the pellets. Furthermore, zirconium doping improved contact between the SE particles. This reduced cracking and increased the relative density of the pellets. This phenomenon can be viewed from various perspectives. Solid-state diffusion plays a significant role in particle formation and interparticle bonding, and therefore, diffusion-based bonding has a vital effect on the mechanical and microstructural properties of materials [42–44]. Furthermore, diffusion itself is governed by sintering temperature and time [42–44]. Consequently, the relative density of the Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs decreased when the zirconium content was higher than x = 0.1. The large particle size of the SEs was attributed to the migration of boundaries [32]. Furthermore, the large particle size reduced the contact area between the SE particles, leading to lower density [32,42]. High zirconium doping concentrations induced discontinuous grain growth. According to these results, particle size growth was directly proportional to doping concentration. In addition, the ionic conductivities of the Li1+xTa1–xZrxSiO5 (x = 0–0.3) pellets prepared with different zirconium doping concentrations were determined by performing an EIS analysis.
(a) XRD patterns of Li1+xTa1–xZrxSiO5 (x = 0–0.3) materials with different compositions, (b) graph of the (110) plane, (c) d (110) plane spacing of Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs with different compositions, and (d) EIS curves of Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs with different compositions.
FE-SEM images of Li1+xTa1–xZrxSiO5 materials for (a) x = 0, (b) x = 0.05, (c) x = 0.1, (d) x = 0.15, (e) x = 0.2, and (f) x = 0.3.
Electrochemical characteristics of Li1+xTa1–xZrxSiO5 (x = 0–0.3) SE by optimizing the doping concentration
A Cole–Cole plot of the Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs sintered at 1050°C for 12 h is presented in Fig. 5d. As depicted in this figure, a semicircle similar to that mentioned in the preceding sections was observed. For x = 0 and 0.05, the SEs exhibited low Li-ionic conductivities of 1.78×10–7 and 2.97×10–7 S cm–1, respectively, owing to the lower density, d-spacing (110 plane) value, and lattice parameters (Table 4). As the Zr concentration increased, the total resistance value decreased at first and then increased, reaching a minimum at x = 0.10 (Rtot =11,500 Ω). The total ionic conductivity of Li1+xTa1–xZrxSiO5 (x = 0.1) was 2.01×10–5 S cm–1. For comparison, the ionic conductivity properties of the LTSO SEs are summarized in Table 6. The heterovalent substitution of Zr4+ for Ta5+ could increase the lithium-ion concentration to accommodate charge compensation. In a similar manner, it may reduce the activation energy and, additionally, induce concerted diffusion to improve Li+ ionic conductivity [24,27,28]. Zr doping in the LTSO SEs enhanced their ionic conductivities to up to 10–5 S cm–1, which is two times higher than those of the pristine LTSO SEs (1.78×10–7 S cm–1). This improvement was ascribed to the reduction of pores and cracks, as well as to the denser structure of the Li1.1Ta0.9Zr0.1SiO5 SE. Nevertheless, the conductivity of the Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs decreased as the doping concentration exceeded a certain value (x > 0.1), which was undoubtedly attributed to an increase in the LiTaO3 peak intensity, decreasing density, and abnormal particle size growth (Fig. 6) triggered by the generation of a new ZrSiO4 phase (Fig. 5a).
Ionic conductivities and relative densities of Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs having different compositions
The ionic conductivities and relative densities of the Li1+xTa1–xZrxSiO5 (x = 0–0.3) SEs are summarized in Table 6. Compared to the pristine LTSO SEs (relative density of 79.5%), the relative densities of the Zr-doped LTSO SEs were higher (in excess of 90%), which first increased and then decreased with increasing doping concentration. It is well established that zirconium doping of LTSO SEs improves their conductivity and mechanical strength, in addition to reducing their resistance. In addition, it was concluded that sintering temperature (1050°C), time (12 h), and zirconium substitution (x = 0.1) conditions increased the relative density and reduced cracking, which helped improve the ionic conductivities of the doped LTSO SEs by 250 times compared to those of the pristine LTSO SEs. Nevertheless, the ion conductivities of the doped LTSO SEs were lower than the theoretical value of 6.4×10–3 S cm–1 calculated through simulation. This was attributed to surface cracks, crystalline collapse, presence of LiTaO3, and changes in particle size during the sintering process. Therefore, the optimization of process variables such as sintering temperature and time, substitution composition, substitution element changes, pressure required for pellet production, particle size, and density should be explored in future studies.
Fig. 7 shows the Galvanostatic charge-discharge, Cyclic voltammetry, and Charge-discharge performance of the LTZSO SE cells. The working current density of the cell was 0.1 mA cm–2 for the corresponding plating and stripping process and the duration time was 30 min. Fig. 7a confirms there was no fluctuations during the cell performance and the electrolyte is stable up to 100 cycles. We added liquid electrolyte at the interface to improve the contact between electrolyte the Li metal electrode. The CV diagram (Fig. 7b) shows clear oxidation and reduction peaks corresponding to lithium intercalation and de-intercalation in the NCM811 cathode. After the first cycle, the oxidation and reduction potential of the cell was shifted to lower potential due to the activation process. From Fig. 7c–d, the initial charge and discharge capacity of the cell were 123 and 99 mAh g–1, respectively. Subsequently after the first cycle, the discharge capacity of the cell was increased due to the activation effect in the electrode/electrolyte interface. Afterwards 15 cycles, the cell provided a discharge capacity of 150 mAh g–1. Moreover, we are working to improve the ionic conductivity, capacity and cell performance.
(a) Galvanostatic charge-discharge curve of the Li//LE-LTZSO-LE//Li symmetric cell with current density 0.1 mA cm–2, (b) Cyclic voltammetry at 0.1 mV s–1 scan rate, (c) Charge-Discharge curve of the cell for selected cycles and (d) Cycle life of the Li//LE-SE-LE//NCM811 cell at 0.2 C rate and all the electrochemical studies were done at 25°C. The NCM811 electrode was prepared by a slurry casting method and the loading mass of the active material was 2–2.5 mg cm-2.
CONCLUSIONS
We successfully synthesized Li1+xTa1–xZrxSiO5 SE by using ball milling and optimized the synthesis by adjusting the sintering temperatures, sintering time, and doping composition. In XRD results of the prepared electrolytes, a few LiTaO3 impurity peaks were observed at all sintering temperatures. In terms of total Li-ion conductivity, the highest value of 7.76×10–8 S cm–1 was achieved by the sample sintered at 1050°C for 6 h. The grain boundary resistance of the SEs decreased at first and then increased as the sintering time exceeded 12 h. Excessive sintering time led to structural instability owing to particle expansion and crystalline collapse, and the optimal sintering time of the LTSO SEs was 12 h. Zr doping played a crucial role during sintering of the packed calcined powders and in stabilizing the SE structure. After sintering, Zr4+ was successfully incorporated into the Li1+xTa1–xZrxSiO5 phase. In terms of optimization of the substitution composition, for x = 0.3, the crystalline phase collapsed, and the proportion of LiTaO3 impurities increased; additionally, a ZrSiO4 peak was generated. After Zr incorporation, the resistance value of the Li1.1Ta0.9Zr0.1SiO5 SE decreased, and it exhibited a high Li-ion conductivity of 2.01×10–5 S cm–1 with a relative density of 90.7%. To further improve the ionic conductivity of the SE, process variables such as sintering conditions, dopant composition, pellet production pressure, particle size, and density should be optimized.
Acknowledgements
This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korean government (Ministry of Science, ICT & Future Planning) (No. RS-2023-00208361). VA acknowledges financial support from the Anusandhan National Research Foundation (ANRF), Govt. of India, through Swarnajayanti Fellowship (SB/SJF/2020-21/12).