General Strategy for Integrated SnO2/Metal Oxides as Biactive Lithium-Ion Battery Anodes with Ultralong Cycling Life.

Integration of bicomponents into a greater object or assemblage is a new avenue to acquire multifunctionality for metal oxide-based anodes for lithium-ion batteries (LIBs). Herein, we report a versatile means by which precursors serve as self-sacrificing templates to form architectures of SnO2 phase and other metal oxides. The vital challenge is the determination of appropriate synthetic system that can benefit the formation of respective precursors in a structure or single-source precursors of tin and other metal species. In the current work, by the aids of synergy action between l-proline and ethylene glycol (EG), precursors containing two metal ions are generally fabricated. Adequate flexibility of the present method has been achieved for SnO2/M x O y hierarchical hybrids, including Mn2O3, Co3O4, NiO, and Zn2SnO4, by calcination of their corresponding SnMn, SnCo, SnNi, and SnZn precursors, respectively. When evaluated as anode materials for LIBs, the obtained SnO2/Mn2O3 homogeneous hybrids, as expected, show higher specific capacity and ultralong cycling stability, gaining a reversible specific capacity of 610.3 mA h g-1 after 600 cycles with only decay of 0.29 mA h g-1 per cycle at 1 A g-1 and 487 mA h g-1 after 1001 cycles at a high current density of 2 A g-1.

Introduction

As one of the excellent energy storage carriers, lithium-ion batteries (LIBs) have drawn extensive attention in the past decades because of the broad range of applications in mobile electric devices, hybrid electric vehicles, and smart grids.[1−6] To fulfill the higher requirements of these applications, further improvements in terms of specific capacity and cycling life should be initiated. Because of their intrinsic limitations in performance, bulk electrode materials fail to completely meet the ever-growing demands. It has been widely demonstrated that nanosized anode/cathode materials enable better rate capability and capacity via reducing diffusion distances of lithium ions and augmenting the contact area between the electrolyte and the electrode. In addition, the commercial graphite anode exhibits a low theoretical capacity (372 mA h g–1),[7,8] which cannot keep pace with the increasing need for energy density of those apparatus with LIB as a power source. In response, metal oxides with higher theoretical capacity (e.g., SnO2,[9,10] Co3O4,[11] MnO2,[12−14] Fe2O3,[15] and ZnCo2O4[16]) have been emerging as attractive anode candidates for LIBs. However, they usually show poor cycling stability and rate capability resulting from low conductivity and remarkable volume change during the charge/discharge process.[9−19] To solve the above issues, it is a useful strategy to assemble two kinds of metal oxides with nanosize into hierarchical composites.[20−22]

Recently, such bicomponent hybrids featured with the complex heterostructure have been developed and designed to serve as active materials for LIBs, verifying that these hybrid anodes are furnished with better lithium storage property (higher specific capacity, better cycling stability, etc.) than any corresponding counterparts.[20−22] From the viewpoint of structural modeling, one representative pattern is that two components heterogeneously dispersed in the structure and can be clearly differentiated due to their definite boundary between each other, such as ZnO@α-Co(OH)2 core–shell structures featured by ZnO spheres as core and α-Co(OH)2 nanosheets as shell[23] and α-Fe2O3/SnO2 heterostructures[24,25] consisting of SnO2 nanowire stem with sixfold α-Fe2O3 nanorod branches and SnO2 nanosheet base growing with branched Fe2O3 nanorods. Two steps are versatilely applied to sequentially generate the two components or precursors into such hybrid architectures with one material as matrix for the growth of another. Another appealing motif is the homogeneous dispersion of components in the entire heterostructure.[26−28] The Yuan group reported the ZnO/ZnFe2O4 submicrocubes constructed with well-dispersed subunits of two oxides by calcination of Prussian blue analogue of Zn3[Fe(CN)6]2 cubes, demonstrating better electrochemical Li storage performance.[26] CoO/CoFe2O4 nanocomposites had been derived from CoFe-layered double hydroxides, exhibiting a tunable lithium storage property by changing the ratio of the component metal oxides.[27] In the connection, one of the challenging issues is to obtain single-resource precursors involving two active metals. Despite great efforts, the effective and general synthesis was largely prohibited by the difficulty of epitaxial growth of the second kind of materials or the limited range of such single sources of bimetal precursors as Prussian blue. Moreover, the aforementioned composites cannot tolerate long cycling at higher current density, probably pertaining to the lack of proper conjoint components. Therefore, efficient coupling of bimetals in a versatile synthetic system should be initiated immediately targeted at high-performance anode materials for LIBs.

Motivated by the advantages of bioxide composites, the general one-pot synthesis of such hybrids was realized in our work by picking a suitable couple of solvent and ligand. This was the first time to enable the architectures of nanosized SnO2 phase and other oxides including Mn2O3, Co3O4, NiO, and Zn2SnO4. By rationally picking up the powerful linker and solvent to conjugate tin and other metal species, hierarchical structures composed of single- or double-source precursors were obtained depending on the intrinsic property of metal ions, which played the key role in the synthesis of bioxide composites. Particularly, SnO2/Mn2O3 hierarchial hybrids (HHs) consisted of nanosheet networks attached to the nanorod stem. When SnO2/Mn2O3 HHs worked as the anode materials for LIBs, a reversible capacity of 487 mA h g–1 was maintained even after a long 1000 cycles at a high current density of 2 A g–1, which was much better than SnO2/Mn2O3 submicrorods and pristine Mn2O3. The origin of performance elevation could be attributed to the synergistic effect of nanohybrid, structural interface, and void among the nanosheets. This strategy displays adequate versatility for architectures of advanced oxide-based hybrid anodes for LIBs and supercapacitors.

Results and Discussion

The overall morphology and structure are investigated by the field emission scanning electron microscopy (FESEM) technique. Figure a demonstrates a high yield of the SnMn precursor. To be more detailed (Figure b,c), some assemblages bind to the submicrorod stem similar to agaric growing on trees. Thereinto, submicrorods are about 5–10 μm in length. Moreover, these assemblages are assembled with interconnected nanosheets, leaving significant voids. According to the thermogravimetric analysis (TGA) result (Figure S1 in the Supporting Information), the precursor was annealed above 500 °C to transform completely into SnO2/Mn2O3 HHs. As shown in Figure d–f, the morphology and structure of SnO2/Mn2O3 HHs inherited well from the paternal template.

Figure 1

FESEM images at different resolutions of the SnMn precursor (a–c) and SnO2/Mn2O3 HHs (d,e). Scale bar: (a) 10 μm, (b) 1 μm, (c) 500 nm, (d) 10 μm, (e) 1 μm, and (f) 500 nm.

The transmission electron microscopy (TEM) technique was applied to visualize the detailed structure information. In Figure a,b, for SnMn precursor, it is clearly exhibited that the thin nanosheets constitute flowerlike clusters which are grown on the submicrorod stem. The TEM images of SnO2/Mn2O3 HHs are displayed in Figure c,d, demonstrating the structural feature inherent from their precursors. The high-resolution TEM (HRTEM) image of Figure e is used to detect the crystalline nature of SnO2/Mn2O3 HHs. Furthermore, the faster Fourier transformation electron diffraction (FFT-ED) pattern was recorded from the square area of Figure e, exhibiting the well-defined diffraction spots. After calculation and analysis, the sets of diffractions are indexed to be (020) and (211) of cubic Mn2O3 with lattice spacings of 0.47 and 0.39 nm, respectively. Another HRTEM image in Figure f was also recorded from SnO2/Mn2O3 HHs. The FFT-ED pattern from the marked square can be indexed to be two sets of diffraction spots, one of which is corresponding to cubic Mn2O3 labeled by a parallelogram and another spot marked by a circle is calculated to be related with the (101) plane of tetragonal SnO2. At the same time, the elemental configuration is sharply visualized by a scanning transmission electron microscopy (STEM) instrument, where the mapping images in Figure g describe the distribution of Mn and Sn mainly focusing on assemblages and submicrorod stems, respectively. The phase of the annealed sample was confirmed by the X-ray diffraction (XRD) pattern in Figure a. They can be well-indexed to be tetragonal SnO2 (JCPDS no. 41-1445) and cubic Mn2O3 (JCPDS no. 41-1442) marked in the pattern. There is no peak for impurity, indicating the high purity of the product.

Figure 2

TEM images of (a,b) SnMn precursor and (c,d) SnO2/Mn2O3 HHs; (e,f) HRTEM images from SnO2/Mn2O3 HHs and (g) corresponding elemental mapping images of Mn, Sn, and O for one single SnO2/Mn2O3 HH.

Figure 3

XRD patterns of (a) SnO2/Mn2O3 HHs, (b) SnO2/Co3O4 HHs, (c) SnO2/NiO HHs, and (d) SnO2/Zn2SnO4 HHs.

To verify the multifunctionality of the synthetic system in the present method in regard of incorporating two metal species into the precursors, we used Co(CH3COO)2·4H2O, Ni(CH3COO)2·4H2O, and Zn(CH3COO)2·2H2O to substitute the previous Mn(CH3COO)2·4H2O for study and successfully fabricated their corresponding precursors which were nominated as SnCo, SnNi, and SnZn precursors, respectively. As shown in Figure S2, these precursors were profiled by similar XRD patterns with well-defined peaks, demonstrating good crystallinity. After calcination at 500 or 600 °C, various SnO2/oxides HHs were harvested, which were analyzed via the XRD technique in Figure . Specifically, the XRD pattern in Figure b of SnO2/Co3O4 HHs can be indexed to tetragonal SnO2 (JCPDS 41-1445) and cubic Co3O4 (JCPDS 42-1467). Simultaneously, other hierarchical hybrids (HHs) of SnO2/NiO and SnO2/Zn2SnO4 can effectively be identified by XRD patterns in Figure c,d, namely, tetragonal SnO2 (JCPDS 41-1445), cubic NiO (Fm3̅m, JCPDS 47-1049), and cubic Zn2SnO4 (JCPDS 24-1470). These XRD results confirmed the formation of various SnO2/oxide hybrids, implying the versatility of the present method.

The FESEM image in Figure a captures a panoramic sight of SnO2/Co3O4 HHs, which displays the screw-nut-like structures that are clearly observed from Figure b. It is clear that the caps are attached at both ends, and some discrete nanoparticles bind on the rodlike stem. Furthermore, the microstructure details of SnO2/NiO HHs were also inspected. Figure c remarkably describes the typical one-dimensional structures. On the basis of the observation from the higher-magnification FESEM image in Figure d, several nanorods are preferentially stacked into bundles with a few nanoparticles scattered onto the nanorod surface. When Zn(CH3COO)2·2H2O replaced Mn(CH3COO)2·4H2O in our synthesis, final SnO2/Zn2SnO4 HHs were transformed from the precursor. It is interesting to note that SnO2/Zn2SnO4 HHs still maintain one-dimensional structures but hexagonal prisms, as indicated in Figure a,b. Because of sintering at high temperature, some prisms suffer from cracking pointed by arrows. The mapping description in Figure c gives the elemental distribution of Zn and Sn, implying the uniform dispersion of SnO2 and Zn2SnO4 domains over the whole hexagonal prism.

Figure 4

FESEM images of SnO2/Co3O4 HHs (a,b) and SnO2/NiO HHs (c,d). Scale bar: 1 μm for all panels.

Figure 5

(a,b) FESEM images of SnO2/Zn2SnO4 HHs and (c) TEM image recorded from a typical single SnO2/Zn2SnO4 hexagonal prism and the corresponding elemental mapping images (c) of Sn and Zn. Scale bars: (a,b) 1 μm and (c) 2 μm.

To shed light on the formation mechanism of precursors, some tests and related analyses were done. In the XRD pattern of the SnMn precursor in Figure S2, the appearance of the peak at about 10.8° demonstrates the formation of a glycolate precursor.[29] A control experiment was carried out where K2SnO3·3H2O was not applied with other conditions being kept constant, and finally, flowerlike Mn–glycolate was obtained, as shown in Figure S3. After heat treatment, pristine Mn2O3 was attained and indexed by the XRD pattern in Figure S4. Similarly, if Mn(CH3COO)2·4H2O was not added, no product was harvested. The comparative results implied that Mn was inclined to bond with EG but not with l-proline and had higher coordination ability to bond with EG than Sn. Hence, it is reasonable to speculate that the SnMn precursor was composed of Mn–glycolate and Sn–proline. To obtain more information about the precursor, Fourier transform infrared (FTIR) spectra of l-proline and SnMn precursor were recorded in Figure S5. A remarkable band at 3433 cm–1 in both spectra implies the occurrence of hydroxyl vibration. In the spectrum of the SnMn precursor, the bands at 2934, 2895, and 2837 cm–1 are related to the asymmetric stretching vibration of N–H and C–H groups.[30] It should be noted that the C=O stretch band for the carboxyl group of l-proline appeared at 1623 cm–1 and disappeared in the spectrum of the precursor.[31,32] There are two new bands at 1633 and 1462 ascribed to the asymmetric νas(COO–) and symmetric νs(COO–) stretching of carboxyl groups, respectively. According to the wavenumber separation value (D) between the νas(COO–) and νs(COO–) bands, the type of interaction between the carboxylate head and metal ions can be distinguished as monodentate, bidentate bridging, bidentate chelating, or ionic interaction.[32−34] Here, the D value is 171 cm–1, suggesting that the carboxyl groups play a role in bidentate bridging.[32−34] Thereafter, the Sn–proline complex is described as Sn–O–C(R)–O–Sn in which l-proline bridges two Sn atoms at each end. Some unsaturated Mn atoms in Mn–glycolate and Sn atoms in Sn–proline would bond with the carboxyl of Sn–proline and hydroxyl of Mn–glycolate, which rendered the hybridization of SnMn precursor HHs. When Zn(CH3COO)2·2H2O was used in the synthetic system, the as-attained SnZn precursor exhibited well-shaped hexagonal prisms. Moreover, the characteristic XRD peak corresponding to the metal glycolate disappears. Correspondingly, the SnZn precursor is inferred to be a single-source analogue of the SnZn–proline complex. After calcination at high temperature, it completely transformed into SnO2 and ZnSn2O4, similar to the reported Prussian blue-structured Zn3[Fe(CN)6]2 as a template to fabricate ZnO and ZnFe2O4 components.[26] These precursors, including SnMn, SnCo, SnNi, and SnZn precursors, showed prominently different coordination properties and crystallization behaviors to get different structure bimetal oxides, which mainly originated from the different intrinsic properties of cations, that is, Mn2+, Co2+, Ni2+, and Zn2+. A more comprehensive and extensive investigation into the underlying fundamentals for such distinctions with changing metal cations is undergoing. From the above discussion, the significance of synthetic system was highlighted. The option of appropriate system can moderate the hybridization fashion and determine the realization of final hybridization.

To prove the superiority of such hierarchical structures as the anode materials, another control sample of SnO2/Mn2O3 submicrorods was attained from its precursor via calcination. As can be seen in Figure S6a, the FESEM image of the precursor of SnO2/Mn2O3 submicrorods describes a rodlike structure with a length of 2–8 μm and a diameter of hundred nanometers to several micrometers. Moreover, this morphology was maintained after annealing, as demonstrated in Figure S6b. The XRD pattern of SnO2/Mn2O3 submicrorods in Figure S7 offers a similar profile, which is determined to be tetragonal SnO2 (JCPDS no. 41-1445) and cubic Mn2O3 (JCPDS no. 41-1442). Further confirmation of chemical composition was given by the STEM mapping results (Figure S6c), evidently demonstrating the homogeneous distribution of Sn, Mn, and O throughout the rodlike structure.

When tested as an anode material for LIBs, the electrochemical behaviors of several samples were characterized. As shown in Figure a, both SnO2/Mn2O3 HHs and submicrorods exhibit good rate capability. At the low current densities of 0.2, 0.5, and 1 A g–1, the capacity of SnO2/Mn2O3 HHs is a little higher than that of the submicrorods. However, when the current density increases to 2 and 5 A g–1, both samples offer a similar reversible capacity. As the rate is back to 0.5 A g–1 in Figure S8, the specific capacity of SnO2/Mn2O3 HHs is 859.2 mA h g–1 and is comparable with that of 814 mA h g–1 at the initial current density of 0.5 A g–1, indicating a good rate capacity of the active material. The reversible capacity can be still maintained as high as 839.9 mA h g–1 after 260 cycles. However, SnO2/Mn2O3 submicrorods rapidly degrade to about 400 mA h g–1 just at 140th cycle.

Figure 6

(a) Rate capability and (b) cycling performance of SnO2/Mn2O3 HHs and submicrorods at a current density of 0.5 A g–1 and (c,d) cycling performance of SnO2/Mn2O3 HHs at current densities of 1 and 2 A g–1, respectively.

Figure S9 shows the cyclic voltammetry (CV) curves of SnO2/Mn2O3 HHs at a scan rate of 0.1 mV s–1 in the potential range of 0.01–3.0 V. In the first cathodic process, a weak and broad cathodic peak was located at around 1.26 V but disappeared in the successive cycles, which can be attributed to the reduction of Mn3+ to Mn2+.[35,36] The main cathodic peak at around 0.96 V is ascribed to the reduction of SnO2 to Sn and Li2O.[37] The last main peak located at 0.24 V is associated with the reduction of Mn2+ to Mn0 and the alloying reaction of Sn and Li+ (LiSn).[35,36,38,39] In the anodic process, the anodic peak at 0.56 V is for the oxidation of LiSn alloy to Sn.[38] The anodic peak at 1.27 V is ascribed to the oxidation of Mn0 to Mn2+ and Sn0 partially back to SnO2.[40,41] In the second scan, the cathodic peak at 0.96 V weakens and then vanishes subsequently, suggesting an irreversible transition of SnO2 to Sn.[37,41] Other two peaks at ∼0.44 and 0.06 V are attributed to the reduction of Mn2+ to Mn0 and the formation of an LiSn alloy, respectively.[35,36,39]

Here, the cycling performance of the as-prepared samples at different current densities was also studied. First, at the current density of 0.5 A g–1 in Figure b, the initial discharge capacity and charge capacity of SnO2/Mn2O3 HHs are 1598.7 and 1021.4 mA h g–1, respectively. The significant capacity loss of 577.3 mA h g–1 is generally considered to be caused by the formation of a solid electrolyte interphase (SEI) layer on the surface of anode materials in the first discharge process.[42−47] Subsequently, the capacity profile undergoes declination followed by inclination, which is familiarly observed for transition metal oxides. This phenomenon generally results from the reversible formation of the gel-like polymeric layer.[48−53] This process was recognized as the activation stage. Even after 300 cycles, SnO2/Mn2O3 HHs could maintain stable capacity, delivering 867.8 mA h g–1, whereas SnO2/Mn2O3 submicrorods show a trend of capacity down all the way and just recover 428 mA h g–1 at the 140th cycle. When the cycling performance of SnO2/Mn2O3 HHs is tested at the current density of 1 A g–1 (Figure c), the first discharge and charge capacities are 1734.9 and 1126 mA h g–1, respectively, with a high initial coulombic efficiency of 65%. The irreversible capacity can be attributed to the decomposition of the electrolyte to form a SEI layer on the surface of the electrode.[42−47] From the sixth cycle onward, the electrode shows an excellent stable capacity, and the reversible capacity of SnO2/Mn2O3 HHs is 610.3 mA h g–1 after 600 cycles with only a slight decay of 0.29 mA h g–1 per cycle. By comparison, as displayed in Figure S10a, the first discharge and charge capacities of SnO2/Mn2O3 submicrorods and pristine Mn2O3 are 1519.5 and 1028.5 as well as 933.7 and 526.1 mA h g–1 at a current density of 1 A g–1. After lithium storage cycling, the reversible capacity of both samples degrades to 192.6 at 250th cycle and 290.4 mA h g–1 at 100th cycle. When the current density increases to 2 A g–1 (Figure d), SnO2/Mn2O3 HHs can still maintain a good cycling performance, which additionally suggests a remarkable rapid charging ability. It delivers a reversible capacity as high as 487 mA h g–1 after 1001 cycles. The other two electrodes made of SnO2/Mn2O3 submicrorods and pristine Mn2O3 cannot endure long cycling at such a high current density. From Figure S10, the reversible capacities of the former electrode are only 192.6 and 115 mA h g–1 after 250 cycles at the current densities of 1 and 2 A g–1, respectively, recovering about 12.7 and 7.6% from the first discharge capacity. In addition, the latter one experiences 100 lithiation/delithiation cycles, and the charge capacities reduce to 290.4 and 202.7 mA h g–1 at the current densities of 1 and 2 A g–1, respectively.

On the basis of the above data and analyses, it was well-concluded that SnO2/Mn2O3 HHs exhibited a superior rate capability and cycling life to SnO2/Mn2O3 submicrorods and pristine Mn2O3. Through the structural and compositional comparison among them, the superiority of SnO2/Mn2O3 HHs can be analyzed as follows. The nanosheet-interconnecting network could raise the surface area of the electrode materials, favorably increasing the contact area between the electrode and the electrolyte and allowing the full occurrence of electrochemical reactions, demonstrated a better performance of SnO2/Mn2O3 HHs than the submicrorods. What is more, the structural voids benefited buffering the volumetric variation of the active materials during the repeated discharge/charge process. Last, the synergetic effect of SnO2 and Mn2O3 was another dominating reason to enhance the electrochemical performance of the composite,[54] which was verified by the worse lithium storage property of the single-phase Mn2O3.

Conclusions

In summary, the general fabrication of integrated SnO2/metal oxide HHs has been devised based on the template-directing route, including Mn2O3, Co3O4, NiO, and Zn2SnO4 as the oxide phase. Their corresponding precursors containing tin and another metal were successfully designed by taking advantage of the synergistic role of l-proline and EG. By analyzing the phase and bond information via XRD and FTIR, single- or double-source precursors could be engineered depending on the coordination and/or intrinsic property of different metal species. In principle, the synthetic strategy is applicable to the integration of other metal oxides. Thereinto, SnO2/Mn2O3 HHs were featured as architectural hybrids with flowerlike clusters anchoring on the rod stem. In accordance with our present work, SnO2/Mn2O3 HHs as the anode material for LIBs showed the enhanced storage capacity and cycling stability. The superior electrochemical performance of SnO2/Mn2O3 HHs could be attributed to their special nanohybrid hierarchical architecture. It proved once more that the reasonable design of nanoarchitectures showed an encouraging avenue to improve the performance of anode materials for LIBs.

Materials and Methods

Synthesis of SnO2/Mn2O3 HHs

All chemical reagents were used as received without any further purification. The synthesis of the SnO2/Mn2O3 hybrid precursor (SnMn precursor) was conducted by a simple refluxing method. In a typical synthesis, 1 mmol Mn(CH3COO)2·4H2O, 0.5 mmol K2SnO3·3H2O, and 1 mmol l-proline were dispersed into 40 mL of EG and oil-bathed in a one-neck capped flask connected with a condenser at 170 °C for 6 h under continuous stirring. After being cooled to room temperature naturally, the product was centrifuged, washed with ethanol several times, and then dried at 60 °C overnight to obtain brown powder. Subsequently, the precursor was calcined at 500 °C for 5 h in an air atmosphere with the heating rate of 2 °C/min. Finally, the dark brown SnO2/Mn2O3 HHs were obtained.

Synthesis of SnO2/Co3O4 HHs

The synthetic procedures were similar to those of SnO2/Mn2O3 HHs except that Mn(CH3COO)2·4H2O was replaced with 1 mmol Co(CH3COO)2·4H2O in the experiment.

Synthesis of SnO2/NiO HHs

The synthetic procedures were similar to those of SnO2/Mn2O3 HHs except that Mn(CH3COO)2·4H2O was replaced with 1 mmol Ni(CH3COO)2·4H2O in the experiment.

Synthesis of SnO2/Zn2SnO4 HHs

The synthetic procedures were similar to those of SnO2/Mn2O3 HHs except that Mn(CH3COO)2·4H2O was replaced with 0.5 mmol Zn(CH3COO)2·2H2O in the experiment, and the corresponding precursor was annealed at 600 °C.

Synthesis of Mn2O3 Nanosheets

The synthetic procedures were similar to those of SnO2/Mn2O3 HHs except that K2SnO3·3H2O was not applied in the experiment.

Synthesis of SnO2/Mn2O3 Submicrorods

The synthetic procedures were similar to those of SnO2/Mn2O3 HHs except that the flask was opened and not connected with the condenser when oil-bathing the reagents.

Characterization of Materials

The phase detection of products was performed by powder XRD (Philips X’Pert Pro Super diffractometer, Cu Kα radiation, λ = 1.54178 Å). TEM and FESEM images were recorded with a transmission electron microscope (JEOL, JEM-1011) and a field emission scanning electron microscope (JEOL, JSM-6700F), respectively. HRTEM images were recorded using an electron microscope (JEM-2100F, accelerating voltage: 200 kV) coupled with an X-ray energy-dispersive spectroscopy (EDX) instrument. TGA (PerkinElmer Diamond TG/DTA apparatus) was conducted at a heating rate of 10 °C/min in flowing air. FTIR spectra were recorded on a Bruker EQUINOX55 spectrometer with a potassium bromide pellet as a control.

Electrochemical Characterization

The working electrode was made of 70 wt % active material, 20 wt % conductive material (acetylene black), and 10 wt % binder (carboxymethylcellulose sodium). The above materials were mixed with water by grinding to form a slurry, which was then pasted on the pure copper foil and dried at 80 °C in vacuum. The obtained foil was cut into a disk with a diameter of 12 mm, and the mass density of the active material was about 1 mg cm–2. The electrochemical measurements were performed using CR2032 coin cells in the voltage range of 3.0–0.01 V. The cells were assembled in an argon-filled glovebox with the lithium foil as the counter electrode and the Celgard 2400 membrane as the separator. The electrolyte was a mixture of 1 M LiPF6 in ethylene carbonate, dimethyl carbonate, and diethyl carbonate in a volume ratio of 1:1:1 purchased from Samsung Chemical Corporation.