Ni-CeO2 Heterostructures in Li-S Batteries: A Balancing Act between Adsorption and Catalytic Conversion of Polysulfide.

Lithium-sulfur (Li-S) batteries have attracted considerable attention over the last two decades because of a high energy density and low cost. However, the wide application of Li-S batteries has been severely impeded due to the poor electrical conductivity of S, shuttling effect of soluble lithium polysulfides (LiPSs), and sluggish redox kinetics of S species, especially under high S loading. To address all these issues, a Ni-CeO2 heterostructure-doped carbon nanofiber (Ni-CeO2 -CNF) is developed as an S host that combines the strong adsorption with the high catalytic activity and the good electrical conductivity, where the LiPSs anchored on the heterostructure surface can directly gain electrons from the current collector and realize a fast conversion between S8 and Li2 S. Therefore, Li-S batteries with S@Ni-CeO2 -CNF cathodes exhibit superior long-term cycling stability, with a capacity decay of 0.046% per cycle over 1000 cycles, even at 2 C. Noteworthy, under a sulfur loading up to 6 mg cm-2 , a high reversible areal capacity of 5.3 mAh cm-2 can be achieved after 50 cycles at 0.1 C. The heterostructure-modified S cathode effectively reconciles the thermodynamic and kinetic characteristics of LiPSs for adsorption and conversion, furthering the development of high-performance Li-S batteries.

Introduction

Lithium–sulfur (Li–S) batteries are considered ideal candidates for next‐generation energy storage devices due to a high theoretical specific capacity (1675 mAh g−1) and energy density (2600 Wh kg−1).[ ] Furthermore, interest in Li–S batteries has been increasing because of the low cost, natural abundance, nontoxicity, and environmental friendliness of sulfur (S).[ ] However, commercial application of these batteries has been severely hindered by several intrinsic characteristics, that is, the insulating property of sulfur and its discharge product Li2S,[ ] large volume expansion (≈80%) from S8 to Li2S during the discharge process,[ ] sluggish redox kinetics of the reduction of S8 or oxidation of Li2S,[ ] and the shuttle effect of soluble lithium polysulfides (LiPSs, Li2S , 2 ≤ n ≤ 8) intermediates,[ ] leading to the low utilization of sulfur active materials and consequently poor cycling performance of Li–S batteries.

To date, strenuous efforts have been made to address the abovementioned issues.[ ] Various carbon materials (e.g., carbon nanotubes,[ ] graphene,[ ] porous carbon,[ ] and carbon nanofibers[ ]) have been widely used as support materials to improve the electrical conductivity of S cathodes and confine S.[ , , ] However, the weak physical interaction between nonpolar carbon materials and LiPSs results in ineffective suppression of the shuttle effect of LiPSs by carbon materials during electrochemical cycling, which cause LiPSs to dissolve in the electrolyte.[ , ] To enhance chemical adsorption with LiPSs, polar materials (including various metal particles,[ , ] oxides,[ , , ] sulfides,[ , ] nitrides,[ ] and carbides[ ]) have been mixed/doped with carbon materials as sulfur hosts for Li–S batteries. Experimental and theoretical studies have shown that polar oxides can effectively immobilize LiPSs by strong chemisorption.[ ] Many polar oxides, such as SiO2,[ ] SnO2 [ ] MnO2,[ ] CeO2,[ , ] TiO2,[ , ] MgO,[ ] and VO2,[ ] have been introduced into cathodes to effectively adsorb LiPSs by chemical interactions. However, most of the metal oxides have poor electrical conductivities and cannot directly gain electrons from a current collector, such that the LiPSs anchored on the carbon surface cannot fully participate in electrochemical reactions. Thus, metal oxides are generally composed/doped with carbon materials as S carriers for cathodes in Li–S batteries. LiPSs can thus diffuse to the conductive carbon surface to complete the reaction.[ , , , , ] Unfortunately, carbon materials exhibit weak catalytic activity in accelerating the redox kinetics of S8 and Li2S. The strong adsorption and weak conversion of LiPSs results in insufficient utilization of S and poor cycling performance. Transition metals, such as Fe,[ ] Co,[ ] and Ni,[ , ] can provide the high electrical conductivity and good catalytic conversion in Li–S batteries, but the nonpolar metal particles cannot effectively anchor LiPSs and suppress the shuttle effect during cycling.

In this work, we innovatively present a Ni‐CeO2 heterostructure‐doped carbon nanofiber (Ni‐CeO2‐CNF) as an S host material for Li–S batteries (Figure  ). Experimental measurements and theoretical calculations show that the Ni–CeO2 heterostructure combines the strong adsorption with excellent catalytic activity and good electric conductivity, where the LiPSs anchored on the heterostructure can directly gain electrons from the current collector and realize a fast conversion between S8 and Li2S (Figure 1b). In the absence of catalysts, the S cathode with CeO2‐doped carbon nanofibers (S@CeO2‐CNF) has strong adsorption for LiPSs but achieves slow conversion of LiPSs (Figure 1c). By sharp contrast, carbon nanofibers (CNFs) without CeO2 and Ni catalysts form a long‐range conductive network but have weak adsorption and achieve slow conversion of LiPSs, leading to poor cycling performance (Figure 1d). As a result, in the presence of Ni‐CeO2‐CNF, the sulfur cathode exhibits both excellent long‐term cycling stability (a low decay rate of 0.046% per cycle during 1000 cycles at 2 C) and a high‐rate capability (553.8 mAh g−1 even at a 3 C rate). Meanwhile, under a sulfur loading up to 6 mg cm−2, a high reversible areal capacity of 5.3 mAh cm−2 can be reached after 50 cycles at 0.1 C. Thus, strong adsorption and fast redox kinetics of LiPSs are regulated in S cathodes, promoting the practical application of Li–S batteries.

Figure 1

a) Schematic of the synthesis of S@Ni‐CeO2‐CNF and the operational principles of b) Ni‐CeO2‐CNF, c) CeO2‐CNF, and d) CNF in Li‐S batteries.

Results and Discussions

SEM images of CNFs, CeO2‐CNFs, and Ni‐CeO2‐CNFs obtained after high‐temperature annealing (900 °C) present a crosslinked nanofiber morphology, where each fiber appears virtually uniform with a diameter of ≈200 nm (Figure  ). A TEM image of CNFs (Figure 2b) shows no nanoparticles on the surface. By contrast, abundant nanoparticles uniformly decorate the CeO2‐CNF and Ni‐CeO2‐CNF surfaces (Figure 2c,e). Clear lattice fringes of CeO2 can be observed in a high‐resolution transmission electron microscopy (HR‐TEM) image of CeO2‐CNF (Figure 2d). In particular, an HR‐TEM image (Figure 2f) of Ni‐CeO2‐CNF exhibits clear lattice fringes of Ni(111) and CeO2 (111), demonstrating the in situ formation of Ni‐CeO2 heterostructures.[ ] The Ni‐CeO2 heterostructures with homogeneous distributions of Ni, Ce and O elements are further confirmed by EDS elemental mappings (Figure 2g).

Figure 2

SEM images of a) CNF, CeO2‐CNF, and Ni‐CeO2‐CNF; TEM images of b) CNF, c) CeO2‐CNF, and e) Ni‐CeO2‐CNF; HR‐TEM images of d) CeO2‐CNF, and f) Ni‐CeO2‐CNF; g) selected area and the corresponding element mappings of Ni‐CeO2‐CNF.

X‐ray diffraction (XRD) patterns (Figure S4, Supporting Information) of CNF, CeO2‐CNF, and Ni‐CeO2‐CNF exhibit relatively intense peaks at 25.5° (JCPDS 26–1076) and 44.5° (JCPDS 50–1086), which are in good agreement with those of graphitized carbon materials.[ ] To further evaluate the graphitic degree of the samples, Raman spectra were analyzed (Figure S5, Supporting Information). There are two typical characteristic peaks at 1350 cm−1 (D band) and 1580 cm−1 (G band) of disordered carbon and ordered graphitic carbon, respectively.[ ] The intensity ratios of the D‐to‐G band (I D/I G), representing the carbon graphitization degree of CNFs, CeO2‐CNFs, and Ni‐CeO2‐CNFs, are all close to 1, indicating similar graphitization of the three materials. The specific surface area and pore structure of CNF (Figure S6, Supporting Information), CeO2‐CNF (Figure S6, Supporting Information) and Ni‐CeO2‐CNF (Figure  ) were characterized using nitrogen adsorption‐desorption isotherms.[ , ] The specific surface areas of CNF, CeO2‐CNF, and Ni‐CeO2‐CNF are 66.7, 232.4, and 159.0 m2 g−1, respectively. The pore volumes of CNF, CeO2‐CNF, and Ni‐CeO2‐CNF are 0.05, 0.14, and 0.13 cm3 g−1, respectively. The BJH pore size distribution curves demonstrate the presence of mesopores in the material structures (2–10 nm). Meanwhile, CNF and CeO2‐CNF presence a small number of micropores, as confirmed by the BJH pore size distribution curves. TGA measurements were used to determine the sulfur contents of the S@CeO2‐CNF and S@Ni‐CeO2‐CNF composites. The weight content of S in the S@CeO2‐CNF and S@Ni‐CeO2‐CNF were determined to be 64 wt% from the change in the respective curve (Figure S7, Supporting Information, Figure 3b). The chemical bonding states of Ni‐CeO2‐CNF were studied by X‐ray photoelectron spectroscopy (XPS). The C 1s spectrum of Ni‐CeO2‐CNF can be fit with four peaks (Figure 3c) corresponding to C═C (284.8 eV), C—C (285.7 eV), C—N (286.5 eV) and C—O (288.5 eV).[ ] The C—N peak was produced by the decomposition of PAN and PVP. Figure 3d shows the XPS spectra of O 1 s, in which the peaks at 531.1, 532.5, and 533.8 eV are attributed to lattice oxygen (O2−), adsorbed oxygen species (O2 2−), and hydroxyl species (OH−), respectively.[ ] The XPS spectra (Figure 3e) exhibits Ce 3d oxidation states at 885.6 and 903.5 eV that are assigned to Ce 3d5/2 and Ce 3d3/2, respectively. The ii and II peaks are assigned to Ce3+. The i, iii, iv, I, III, and IV peaks are related to Ce4+.[ , ] The high‐resolution XPS spectra (Figure 3f) of Ni 2p in Ni‐CeO2‐CNF can be fitted to the zero‐valence state (854.5 and 870.1 eV) and ionic state (856.5 and 872.5 eV) with satellites from the Ni—O species (861.1 eV), verifying the existence of metallic Ni. The C 1s, Ce 3d and Ni 2p (Figure S8, Supporting Information) of S@Ni‐CeO2‐CNF were further analyzed. The results show that Ni‐CeO2‐CNF loaded with S is stable.

Figure 3

a) N2 adsorption‐desorption isotherms of Ni‐CeO2‐CNF (the inset shows the pore size distribution); b) TGA curves of S and S@Ni‐CeO2‐CNF; XPS spectra of c) C 1 s, d) O 1 s, e) Ce 3d, and f) Ni 2p of Ni‐CeO2‐CNF.

To corroborate the chemical interaction between the host matrix and LiPSs, a polysulfide adsorption visualization test was conducted on CNF, CeO2‐CNF, and Ni‐CeO2‐CNF. The addition of CeO2‐CNF and Ni‐CeO2‐CNF to a Li2S6 solution causes the solution color to change from dark yellow to colorless (transparent, inset of Figure  ). Soluble Li2S6 is almost completely adsorbed by CeO2‐CNF and Ni‐CeO2‐CNF, as confirmed by UV–vis results, in which the characteristic peaks of Li2S6 (280 nm) and Li2S4 (410 nm) are absent (Figure 4a).[ ] By contrast, the solution containing CNFs and the blank Li2S6 solution remain yellow, and distinct characteristic peaks of Li2S6 and Li2S4 appear in the UV–vis spectra (Figure 4a). All the above mentioned results indicate that the CeO2‐CNF and Ni‐CeO2‐CNF composite powders have a strong adsorption ability for soluble LiPSs. To determine the catalytic activity of Ni‐CeO2‐CNF toward the reduction and oxidation of LiPSs, cyclic voltammetry (CV) measurements were performed on symmetrical cells with CNF, CeO2‐CNF, and Ni‐CeO2‐CNF as the working and counter electrodes. In Figure 4b, the CV curves of Ni‐CeO2‐CNF show four pronounced redox peaks and significantly higher current‐density peaks than those of CNF and CeO2‐CNF, implying superior electrocatalytic activity for S reaction kinetics.[ ] To analyze the superiority of Ni‐CeO2‐CNF toward Li2S nucleation, potentiostatic discharge curves (Figure 4c) of CNF, CeO2‐CNF, and Ni‐CeO2‐CNF cells were tested at 2.05 V with an Li2S8 solution as the electrolyte. The Ni‐CeO2‐CNF electrode exhibits the earliest and highest current responses (0.252 mA at 2170 s for Ni‐CeO2‐CNF, 0.127 mA at 8185 s for CeO2‐CNF, and 0.068 mA at 20360 s for CNF). The Ni‐CeO2‐CNF electrode also exhibits a higher Li2S precipitation capacity (227.1 mAh g−1) than the CeO2‐CNF (93.6 mAh g−1) and CNF electrodes (75.0 mAh g−1). This result indicates a lower barrier to Li2S formation on the Ni‐CeO2 heterostructures compared to CeO2‐CNF and CNF.[ ] The diffusivity of Li ions was determined from the CV curves of S@CNF (Figure S9), S@CeO2‐CNF (Figure S10, Supporting Information) and S@Ni‐CeO2‐CNF (Figure 4d) cathodes assembled in Li–S batteries. The linear relationship between the peak current density (I peak) and the square root scanning rate (v 0.5) is expressed by the Randles–Sevcik equation: I p = (2.65 × 105)n 1.5 SD Li+ 0.5ΔC Li+ v 0.5, where n denotes the number of transfer charges, S denotes the surface area of the active electrode, D Li+ denotes the diffusion coefficient of Li ions, and ΔC Li+ denotes the concentration of Li ions.[ ] The peak currents at the square root of different scan rates are shown in Figure 4e. The slopes of the S@CNF, S@CeO2‐CNF, and S@Ni‐CeO2‐CNF curves are 2.18, 3.11, and 5.83, respectively, at peak 1; 1.10, 1.40, and 2.28, respectively, at peak 2; and 0.98, 2.27, and 2.98, respectively, at peak 3. The significantly larger slope corresponding to the S@Ni‐CeO2‐CNF compared to those of the other cathodes demonstrates a higher diffusion rate of Li ions and faster kinetics of LiPSs conversion during the redox process. To further investigate the internal resistances of cells with S@CNF, S@CeO2‐CNF, and S@Ni‐CeO2‐CNF cathodes during the discharge/charge process, galvanostatic intermittent titration (GITT) profiles were measured at 0.1 C. The internal resistances to nucleation and activation of Li2S are indicated by the dip depth in the discharging and charging profiles (shown by the arrows in Figure 4f–h).Polarization during the discharge−charge process can be quantified by the relative size of ΔR internal in GITT tests according to the following relation: where ΔV QOCV–CCV is the voltage difference between the points of quasi open‐circuit voltage (QOCV) and closed‐circuit voltage (CCV), and I applied is the applied current, as shown in Figure 4f–h inset. The S@Ni‐CeO2‐CNF battery exhibits smaller ΔR internal values than the other batteries (Figure 4i) between the Li2S nucleation and activation points, indicating that the S@Ni‐CeO2‐CNF electrodes have the lowest internal resistance. All the results of the abovementioned experiments indicate that the Ni‐CeO2 heterostructures have a strong adsorption ability for soluble LiPSs and effectively accelerate the electrochemical reaction kinetics during the redox process.

Figure 4

a) UV–vis spectra and optical images of CNF, CeO2‐CNF, and Ni‐CeO2‐CNF with adsorbed LiPS solutions; b) CV profiles of CNF, CeO2‐CNF, and Ni‐CeO2‐CNF symmetric cells; c) potentiostatic discharge profile at 2.05 V for different electrodes with a Li2S8 catholyte to evaluate the nucleation kinetics of Li2S; d) CV curves of the S@Ni‐CeO2‐CNF electrode at various scanning rates; e) profiles of CV peak currents versus the square root of the scan rates; galvanostatic intermittent titration (GITT) voltage profiles of f) S@Ni‐CeO2‐CNF, g) S@CeO2‐CNF, and h) S@CNF cathode at 0.1 C; i) internal resistances of the S@Ni‐CeO2‐CNF, S@CeO2‐CNF, and S@CNF electrodes relative to the normalized discharge–charge time.

The CV profiles of the cells with S@Ni‐CeO2‐CNF between 1.7 and 2.7 V at a scan rate of 0.05 mV s−1 are clearly shown in Figure  . The two cathodic peaks at 2.328 and 2.071 V correspond to the reduction of S8 to high‐order LiPSs and further reduction of LiPSs to short‐chain solid Li2S2/Li2S. The two anodic peaks are attributed to the reverse process of transforming solid Li2S2/Li2S to LiPSs (2.316 V) and further to S8 (2.369 V).[ ] Compared to S@CNF and S@CeO2‐CNF (Figure 5a), S@Ni‐CeO2‐CNF has a cathodic peak at a higher potential (2.071 V) and anodic peaks at lower potentials (2.316 V), indicating enhanced reaction kinetics. Good overlap of the initial three CV cycles also suggests good electrochemical reversibility for the cell with S@Ni‐CeO2‐CNF (Figure S11, Supporting Information).

Figure 5

Electrochemical performances of the S@CNF, S@CeO2‐CNF, and S@Ni‐CeO2‐CNF cathodes: a) CV profiles; b) galvanostatic discharge/charge profiles for the first cycle at 0.1 C; c) EIS curves; d) rate performances and cycling performance at e) 0.5 C and f) 2 C of S@CNF, S@CeO2‐CNF, and S@Ni‐CeO2‐CNF cathodes; g) cycling performance of S@Ni‐CeO2‐CNF cathode at 0.1 C with high S loadings of 6.0 mg cm−2.

The galvanostatic discharge/charge profiles of batteries with different materials at 0.1 C are compared in Figure 5b. The initial discharge capacity is 1327.6 mAh g−1 for S@Ni‐CeO2‐CNF, 1286.2 mAh g−1 for S@CeO2‐CNF, and 1190.2 mA h g−1 for S@CNF. The discharge/charge plateaus are significantly longer for S@Ni‐CeO2‐CNF than the other materials, resulting in a higher capacity. Notably, the cell with the S@Ni‐CeO2‐CNF cathode shows a stable and flat discharge plateau with the lowest polarization potential of 0.116 V. The electrochemical impedance spectroscopy (EIS) curves of the different batteries after ten cycles are shown in Figure 5c. The semicircle in the high‐frequency region reflects the charge‐transfer resistance at the carbon interface, and the semicircle in the medium‐frequency region indicates the formation of a solid Li2S@Li2S2 film on the carbon interface.[ ] The cell with S@Ni‐CeO2‐CNF produces the smallest semicircle in the Nyquist plot, indicating high transportation of ions and more ideal conductivity, which can also decrease the charge‐transfer resistance.

The cells with S@Ni‐CeO2‐CNF also display the highest rate performance for rates ranging from 0.2 to 3 C (Figure 5d). The discharge capacity of S@Ni‐CeO2‐CNF reaches 1197.9 (0.2 C), 938.6 (0.5 C), 779.1 (1 C), 653.2 (2 C), and 572.8 mAh g−1 (3 C). An average specific capacity of 553.8 mAh g−1 can be achieved even at a high rate of 3 C. As the current rate is returned from 3 to 0.5 C, S@Ni‐CeO2‐CNF retains a high average specific capacity of 813.6 mAh g−1. This result indicates that the Li–S batteries with S@Ni‐CeO2‐CNF have a good rate capability and cycle stability. In Figure S12, Supporting Information, the capacity of the S@Ni‐CeO2‐CNF cathode is initially 1208.7 mAh g−1 and maintained at 919.5 mAh g−1 after 100 cycles at 0.2 C. The capacity of the S@Ni‐CeO2‐CNF cathode is initially 933.4 mAh g−1 and maintained at 783.6 mAh g−1 after 300 cycles at 0.5 C (Figure 5e). By contrast, the capacities of the S@CNF and S@CeO2‐CNF cathodes decay faster (329.0 and 520.8 mAh g−1, respectively) after 300 cycles. Moreover, the S@Ni‐CeO2‐CNF cathode exhibits superior cycling performance with a small capacity decay of 0.046% per cycle during 1000 cycles at 2 C (Figure 5f). To further explore the electrochemical performance of S@Ni‐CeO2‐CNF cathodes, a series of tests were carried out under high S loadings of 3.0, 4.0, and 6.0 mg cm−2. A high areal capacities of 5.3 mAh cm−2 (0.1 C) after 50 cycles can be obtained under a sulfur loading of 6 mg cm−2 (Figure 5g). When S loadings are 3.0 and 4.0 mg cm−2, the initial specific capacities at 0.1 C after the activation process are 1116.4 and 1091.8 mAh g−1, respectively, and the remaining capacities after 80 cycles are 774.1 and 683.8 mAh g−1 (Figure S13, Supporting Information). Therefore, these results illustrate the good electrochemical performances of the S@Ni‐CeO2‐CNF cathode even under a high S loading.

Figure  show the binding energies of Ni, CeO2, and Ni‐CeO2 heterostructures with S8, Li2S8, Li2S6, Li2S4, Li2S2, and Li2S. The Ni‐CeO2 heterostructure exhibits the stronger adsorption ability for soluble LiPSs than Ni and CeO2. Therefore, the Ni‐CeO2 heterostructure can efficiently mitigate the shuttling of LiPSs.[ , ] To explain the improved conversion of LiPSs by the S@Ni‐CeO2‐CNF cathode, we performed DFT calculations for the Gibbs free energies of different possible reactions of LiPSs on the Ni‐CeO2 heterostructure and compared the results to those for similar reactions on Ni and CeO2. During the discharge process, S8 undergoes double reduction with two Li+ ions to form Li2S8. Subsequently, Li2S8 undergoes further reduction and disproportionation with the stepwise formation of Li2S6, Li2S4, Li2S2, and the end product Li2S.[ ] The Gibbs free energies of the abovementioned reactions on the Ni, CeO2 and Ni‐CeO2 heterostructure models are shown in Figure 6c, and the optimized structures of the intermediates on the Ni‐CeO2 heterostructure models are shown in the inset of Figure 6c. After the spontaneous conversion of S8 to Li2S8, the subsequent four steps for the formation of Li2S6, Li2S4, Li2S2, and Li2S on all models are either spontaneous (negative Gibbs free energy) or endothermic (positive Gibbs free energy). During the entire discharge process, the rate‐limiting step has the largest positive Gibbs free energy among all the steps, which is 1.21 eV for Ni, 0.67 eV for CeO2, and 0.38 eV for the Ni‐CeO2 heterostructure. The lowest Gibbs free energy of the rate‐limiting step for the Ni‐CeO2 heterostructure indicates that the reduction of S is thermodynamically more favorable on the Ni‐CeO2 heterostructure than on the Ni and CeO2 substrates. We calculated the density of states (DOS) of the Ni, CeO2 and Ni‐CeO2 heterostructures (Figure 6d). The Ni‐CeO2 heterostructure is considerably more metallic with a higher DOS at the Fermi level compared to CeO2. The high electrical conductivity of the Ni‐CeO2 heterostructure will help the LiPSs on the heterostructure surface directly gain electrons from the current collector, and enhance the rapid electrochemical reactions of S and LiPSs. The results of the abovementioned calculations effectively explain the comprehensively excellent performance of Ni‐CeO2‐CNF for sulfur hosts.

Figure 6

a) Optimized binding geometric configurations and energies of Ni, CeO2 and Ni‐CeO2 heterostructures with S8, Li2S8, Li2S6, Li2S4, Li2S2, and Li2S; b) binding energies of Ni, CeO2, and Ni‐CeO2 heterostructures with S8, Li2S8, Li2S6, Li2S4, Li2S2, and Li2S calculated using DFT; c) energy profiles for the discharging process from S8 to Li2S on the Ni, CeO2, and Ni‐CeO2 heterostructure models (insets show the optimized adsorption conformations of intermediate species on the Ni‐CeO2 heterostructure models); d) total density of states for the Ni and Ni‐CeO2 heterostructures.

Conclusion

In summary, we developed a novel sulfur host material, Ni‐CeO2 heterostructure‐doped carbon nanofibers, that combines strong adsorption with high catalytic activity and good electrical conductivity, striking a balance between adsorption and catalytic conversion of LiPSs. The cross‐linked carbon nanofiber can also enhance electrical conductivity of the cathode and act as a second barrier to block the LiPSs shuttling. Therefore, in the presence of Ni‐CeO2‐CNF, the S cathode exhibits both excellent long‐term cycling stability (a low decay rate of 0.046% per cycle during 1000 cycles at 2 C) and a high‐rate capability (553.8 mAh g−1 at a 3 C rate). Moreover, a high reversible areal capacity of 5.3 mAh cm−2 can be obtained even after 50 cycles at 0.1 C when the S loading is up to 6 mg cm−2. The present study provides an effective strategy for accommodating the thermodynamic and kinetic characteristics of LiPS adsorption and conversion. The material prepared in this study has considerable application potential for use in high‐performance Li‐S batteries. Considering the high reversible areal capacity under high S loading, such the Li‐S batteries show great application potentials in drones, electric vehicles, portable electronics, and medical devices.

Conflict of Interest

The authors declare no conflict of interest.

Supporting Information

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