Huaifang Shang, and Dingguo Xia, Spinel LiMn2O4 integrated with coating and doping by Sn self-segregation, Int. J. Miner. Metall. Mater., 29(2022), No. 5, pp.909-916. https://dx.doi.org/10.1007/s12613-022-2482-8
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In recent years, lithium-ion secondary batteries have played a crucial role in the development of modern society as an important energy storage device in portable electronic products and new energy vehicles [1–4]. Layered cathode materials containing cobalt or nickel are among the most promising cathodes for high-energy lithium-ion batteries because of their high capacity and working potential [5–6]. However, due to the high cost and limited resources of cobalt and nickel, it is necessary to develop cobalt-free and nickel-free cathode materials [7–9]. Lithium manganate (LiMn2O4) spinel is a mainstream commercial cathode material for rechargeable energy storage owing to its stable structure, good safety, high operation voltage, and low cost [10–11]. Batteries with LiMn2O4 as the cathode are widely used in low-speed electric vehicles and portable electronic products for energy storage [12–13]. However, in its electric vehicle application, the cycle stability of the battery with LiMn2O4 as the cathode is unable to meet the actual demand when operating at high temperatures for a long time. Thus, the application of LiMn2O4 in electric vehicle power batteries has been limited. The capacity decay of LiMn2O4 is closely related to manganese dissolution caused by the Jahn–Teller effect and electrolyte decomposition. To address these problems, several techniques have been employed, such as ion doping, surface coating, and adding electrolyte additive [14–15]. Current research work focuses on inhibiting the Jahn–Teller phase transition from cubic to tetragonal [16–17] and preventing manganese dissolution in the electrolyte [18], especially at a high temperature at which the average valence state of Mn is below 3.5 [19–20]. However, in many cases, either the improvement effect is not satisfactory, or the preparation process is too complex.
Considering the convenience of the preparation and avoiding the surface two-phase lattice mismatch caused by improper coating on the LiMn2O4 surface, we report for the first time, LiMn2O4 with Sn-rich coating through Sn surface self-segregation, thus integrating coating and doping. The as-obtained sample delivers superior high initial capacities of 124 mAh·g−1 and 120 mAh·g−1 with the capacity retention of 91.1% and 90.2% at 25°C and 55°C after 50 cycles, respectively. Auger electron spectroscopy (AES), powder X-ray photoelectron spectroscopy (XPS), and near-edge X-ray-absorption fine structure were employed to further reveal the origin of the Sn surface enrichment and the suppression of Jahn–Teller distortion. Sn enrichment on the LiMn2O4 surface stabilizes the LiMn2O4 crystal structure and protects the LiMn2O4 cathode from corrosion induced by organic electrolytes.
First, LiNO3, Mn(NO3)2, and SnCl2 were weighed at a molar ratio of Li/Mn/Sn = 1.03:(2−x):x with x = 0, 0.005, 0.01, 0.02, and 0.05, respectively, and completely dissolved into a common solution of purified water with citric acid as a chelating agent. Then, the dried gel was obtained by rotary evaporation at 80°C. Finally, the dried powder was mixed and heated at 850°C for 24 h in the air to obtain the desired materials.
The X-ray diffraction (XRD) measurements were performed using the Bruker D8-ADVANCE diffractometer (Bruker, Germany) with a Cu Kα radiation source (λ = 0.15406 nm). A scanning electron microscope (SEM; Hitachi, S-4800) and a transmission electron microscope (TEM; TECNAI, F20) were used to observe the surface morphology of materials and identify elemental distribution. Transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HRTEM) were employed to study the microstructure of materials on an F20 transmission electron microscope. Powder XPS (AXIS, ULTRA) was performed to analyze the chemical valence states of elemental Sn and Mn, and the data were calibrated using an indefinite C 1s peak with a fixed value of 284.8 eV. AES (PHI, 700) with 5-keV initial electron energy, 15-nA initial current density, and 9-nm/min sputtering speed was used to determine the element type and material content according to the energy and number of auger electrons emitted by the electron beam laser. The O K-edge X-ray absorption near edge structure (XANES) was measured at beamline 4B7B of the Beijing Synchrotron Radiation Facility. The theoretical simulations of O K-edge X-ray absorption spectroscopy (XAS) were performed within the framework of the multiple-scattering theory [21–22] using the FEFF 8.2 code [23] within the muffin-tin approximation. The FEFF input files were generated using the ATOM package, and the Hedin–Lundqvist exchange correlation potential was selected for the calculations [24]. The occupation of multiple positions by O atoms led to some difficulty in calculating FEFF. XANES technology is a well-known response to the average structural information of the different positions of atomic species. Therefore, during the processing, we first performed the XANES spectroscopy for individual sites, followed by the weighted superposition for all the sites.
Electrochemical measurements were performed with the 2032-type coin cells at 25°C and 55°C, respectively, using a Neware battery tester (5 V, 10 mA). Lithium metal was used as the counter electrode. The cathode was prepared by casting a slurry of 80wt% active oxide, 10wt% poly(vinylidene difluoride) binder in N-methylpyrrolidinone solvent, and 10wt% acetylene black on an Al foil substrate.
Fig. 1(a) shows the XRD patterns of LMO–Snx (0 ≤ x ≤ 0.05) and pristine LiMn2O4. All peaks can be indexed by the well-defined spinel phase with the space group of Fd
Fig. 2 exhibits the SEM images of pristine LiMn2O4 and LiMn2O4 with different amounts of Sn. After heat treatment, the LMO–Snx (x = 0, 0.005, and 0.01) products displayed the morphologies of well-crystallized octahedral particles, approximately 400 nm in size, as shown in Fig. 2(a)–(c). An increase in the Sn amount changed the appearance of octahedral particles from clear to vague, and the surface became rough when the Sn content was larger than 0.01 (Fig. 2(d)–(e)).
Fig. 3 shows the TEM images of LiMn2O4 samples with different Sn amounts. Fig. 3(a) shows the surface of the pure phase LiMn2O4 is relatively smooth, while the samples with different Sn contents show different surface topography. Fig. 3(b) shows the TEM images of the LiMn2O4 sample with Sn content of x = 0.005. Due to the low doping, the particle surface topography seemed discontinuous and uneven. Fig. 3(c) and (d) shows that when the Sn content increases to x = 0.01, a uniform and smooth coating layer with a thickness of about 5 nm appears. This may prevent the corrosion of material surface by the electrolyte, inhibit the dissolution of manganese ions, and improve the cycle stability. When the Sn content is 0.02 or 0.05, the particle surface is covered with a 10–30 nm-thick coating layer, as shown in Fig. 3(e) and (f). This thick coating may decrease the specific capacity of the material owing to its low conductivity. The high-resolution TEM analysis of the sample with x = 0.02 shows that the coating layer is SnO2, characterized with the (110) planes of a lattice spacing of 0.33 nm for SnO2, which is consistent with the SnO2 peaks obtained from XRD [26]. Meanwhile, HRTEM images showed that the crystal plane spacing of the pristine LiMn2O4 is 0.448 nm, corresponding to the (111) planes of the LiMn2O4 spinel phase structure, as shown in Fig. 3(g). The crystal plane spacing changes to 0.424 nm after Sn doping (x = 0.005), indicating that the dopant enters the lattice position and decreases the crystal plane spacing, as shown in Fig. 3(h).
Fig. 4 shows the selected particle analysis locations of LMO–Snx when x = 0.005 (Fig. 4(a)) and x = 0.01 (Fig. 4(c)), and Fig. 4(b) and (d) shows the corresponding analysis results of Sn and Mn using the Auger electron energy spectrum. Sn and Mn show clearly different trends with the increase of detection depth. From the edge of the particles to 10 nm inside, Mn shows an increasing trend, while Sn decreases. The strong Sn signal at the particle edge demonstrates the Sn-rich surface layer of Sn-doped LiMn2O4.
The XPS spectra of Mn and Sn for LMO–Snx were overlaid for comparison. The binding energy of the peak maxima for Mn of LMO–Snx shifted toward a higher energy with x from 0.00 to 0.01 (Fig. 5(a)), suggesting a decreased contribution of the Mn(III) species over the substitution range of Sn; however, the peak shifted to a lower energy when x = 0.02 and 0.05, indicating the increase in Mn(III) content. To obtain further information on the Mn percentage at the surface of LMO–Snx, the Mn 2P3/2 XPS spectra were fitted using the XPSPEAK 4.1 software. The Mn XPS spectra in Fig. 5(b) represent the manganese content, which can be used to calculate the average valence state of Mn between Mn3+ and Mn4+. When the binding energy positions of Mn3+ and Mn4+ are fixed at 642.1 and 643.6 eV, respectively, deconvolution can be conducted accurately [27]. The valence states in the x = 0.005 and 0.01 samples were determined as 3.54 and 3.59, respectively, which are higher than the 3.50 value for pristine LiMn2O4. For the Sn contents of x = 0.02 and 0.05, the valance states were 3.53 and 3.52, respectively. This slight increase in the valence state is related to the substitution of Mn3+ at the octahedral site (16d) by Sn2+ to maintain charge neutrality. The XPS spectra of Sn 3d3/2 and Sn 3d5/2 of LMO–Snx are shown in Fig. 5(c), resulting from the spin-orbit splitting. The peak points of both Sn 3d3/2 and Sn 3d5/2 shifted to a lower binding energy for x = 0.005 and 0.01, indicating that the relative content of Sn2+ is greater than that of Sn4+ at a lower Sn content; however, when Sn content is high, it moves toward a higher energy owing to the SnO2 coating layer formation. The Sn content could be identified more definitely from the fitted profiles shown in Fig. 5(d). When the binding energy positions of Sn2+ and Sn4+ were fixed at 485.75 and 486.75 eV, respectively, deconvolution was performed accurately [28]. The average valence of Mn increased for the Sn contents of x = 0.005 and 0.01, whereas that of Mn decreased for x = 0.02 and 0.05 when the intensity of the SnO2 (110) peak increased. This implies that a small amount of Sn2+ replaces Mn3+ with limited solubility, and the remaining Sn2+ is located outside the LiMn2O4 particles as the ion radius of Sn2+ (0.0930 nm) is larger than that of Mn3+ (0.0785 nm) [29], i.e., the average valence state of Mn decreases when x > 0.01.
To further determine the surface phase composite of LiMn2O4 doped with Sn, the K-edge absorption spectrum of O with TEY mode was measured. Fig. 6(a) shows the K-edge absorption spectrum of O in the sample with the Sn content of 0.01 (labeled as exp) and the calculated spectrum of the potential species, including pure LiMn2O4 (LMO), Li2SnO3 (LSO), SnO, and SnO2. The results showed that the three significant feature peaks of LiMn2O4, pre-edge peaks A1 and A2 and peak B, are consistent with the experimental spectrum when the Sn amount is 0.01. When all Mn is replaced by Sn or the surface coating is regarded as an oxide of Sn, the peak type and peak position of the calculated spectrum line are no longer consistent with the experimental spectrum. Using the fingerprint of the O K-edge XAS, we could easily exclude the possibility of LSO, SnO, and SnO2 presence on the sample surface. As evident, the pre-edge peaks correspond to the transition of O 1s electrons into the hybridized state of the metal 3d with O 2p orbitals. The integral intensity of the pre-edge peak of the O K-edge spectra is related to the total d electron hole number of the hybridized states. Moreover, the A2 peaks are more sensitive to the local environment as compared to the A1 peaks. After Sn doping, there is an enhanced intensity of pre-edge peaks A2, meaning that M–O bond becomes stronger and presents features similar to the experimental spectrum. Combined with the XPS results, it could be inferred that the surface layer contains a small part of Sn2+, changing the valence state of Mn. The calculated spectrum of part Mn replaced by Sn (LMO + S) is consistent with the experiment spectrum. In other words, the surface phase can be determined as LiMn1.99Sn0.01O4.
Fig. 6(b) shows the L edge absorption spectra of Mn. The edge position shows an evident shift to a higher energy after Sn doping (enlarged image in Fig. 6(c), demonstrating an enhanced Mn valence state. However, when the Sn contents are x ≥ 0.02, their peak positions shift slightly toward the lower energy. The reason could be that when the Sn2+ proportion increases, it is easier to form SnO2 rather than replace Mn3+; thus, the average valence state of Mn decreases.
The cyclic performance of the pristine LiMn2O4 and LiMn2O4 doped with Sn at room temperature with metal lithium as the counter electrode is shown in Fig. 7(a). Since Sn ion is not involved in the redox reaction in LiMn2O4, high tin doping will reduce the capacity of the material. When the Sn content is 0.01, the first discharge capacity is 124 mAh·g−1. After 50 cycles, the discharge capacity drops to 113 mAh·g−1 with a retention rate of 91.1%, which is apparently higher than that of the pristine LiMn2O4 [30]. The increase in the specific capacity after Sn doping can be attributed to the reduced disproportionation reaction of Mn3+ and prevents manganese dissolution effectively due to the SnO2 coating layer. Moreover, the enhanced electrochemical performance can be attributed to the increase in conductivity after Sn doping [31]. When the amount of Sn increases to 0.02, the first discharge capacity is 122 mAh·g−1, drops to 110 mAh·g−1 after the subsequent 50 cycles. When the amount of Sn is 0.05, the discharge capacity decreases from 113 to 106 mAh·g−1 with the same electrochemical process as above. This may be because excessive Sn makes it easy to form SnO2 particles on the sample surface, resulting in increased material impedance and reduced specific capacity. Fig. 7(b) shows the cyclic performance of the pristine LiMn2O4 and LiMn2O4 doped with Sn at 55°C. When the doping amount is 0.01, the first discharge capacity at 0.2 C rate was 120 mAh·g−1 with a retention rate of 90.2%. When the doping amount is 0.02, the performance is slightly lower but still higher than that of pure lithium manganate.
Fig. 7(c) demonstrates the application of modified material in practice, graphite as an anode, to measure the electrochemical properties at 55°C. For the Sn content of 0.01, the first discharge capacity is 110 mAh·g−1 at 0.5 C, with 108 mAh·g−1 at 1 C and 105 mAh·g−1 at 2 C, while the first discharge capacities of the pristine LiMn2O4 are 109 mAh·g−1 at 0.5 C, 102 mAh·g−1 at 1 C, and 101 mAh·g−1 at 2 C. After 50 cycles, the discharge capacity of the sample with Sn content of 0.01 is 105 mAh·g−1 with a retention rate of up to 95.4%, while the discharge capacity of the pristine sample is 92 mAh·g−1 with a retention rate of only 84.4%, as shown in Fig. 7(d). The electrochemical performance achieved herein is highly competitive with most reported work on modified LiMn2O4 [32–33]. The high electrochemical performance indicates that the surface-coated LiMn2O4 has better electrochemical cycle stability than the pure phase, especially with the significantly improved high-temperature stability of LiMn2O4. As well-known, the Jahn–Teller effect will aggravate the disproportionation reaction of Mn3+. The resulting Mn2+ ion is easily soluble in the electrolyte and destroys the solid electrolyte interface (SEI) film on the anode, declining the battery performance. Moderate Sn doping can enhance the Mn4+ content in the material and inhibit Jahn–Teller effect. Moreover, the Sn-rich coating formed on the surface of LiMn2O4 plays an important role in protecting the material from corrosion by the electrolyte.
In conclusion, LiSn0.01Mn1.99O4 prepared via sol–gel method showed an exceptionally strong cycle ability, with a remarkable capacity retention of 91.1% and 90.2% at the rate of 0.2 C at 25°C and 55°C as well as high initial capacities of 124 mAh·g−1 and 120 mAh·g−1, respectively. The properties of the material depended closely on the dopant amount, which is attributed to the fact that an appropriate amount of the Sn dopant can improve the valence state of Mn and inhibit the Jahn–Teller distortion effectively. Our work demonstrates an effective strategy for improving the electrochemical performance of spinel cathode materials with the integration of coating and doping by Sn self-segregation, and it highlights the role of segregation and optimal designing of spinel cathode materials for long-life and low-cost lithium batteries.
This work was financially supported by the International Science & Technology Cooperation of China (No. 2019YFE0100200), the National Natural Science Foundation of China (No. 53130202), and the Basic Research Program of Shanxi Province, China (No. 20210302123259).
The authors declare no competing financial interest.
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