Anchoring electron-delocalized CeO₂ on porous carbon for expediting polysulfide kinetics toward high-loading Li–S batteries

Abstract

High-energy-density lithium sulfur batteries with high mass loading are restricted by the depressive electrochemical kinetics of polysulfide conversion. Herein, to enhance catalytic efficiency, abundant electron-delocalized CeO₂ nanoparticles are anchored on the surface of pollen-derived carbon (PC-CeO₂) via one-step carbonization, serving as a sulfur host. In this design, pollen-derived carbon (PC) with a porous structural network enhances the electrical conductivity of the sulfur cathode while alleviating volume expansion and maintaining the stability of the cathode. The strategic incorporation of electron-delocalized CeO₂ nanoparticles is beneficial for the adsorption and catalysis of polysulfides, limiting the shuttle effect of polysulfides and effectively facilitating the electrochemical conversion kinetics. As a result, the fabricated sulfur cathode (PC-CeO₂/S) exhibits excellent electrochemical stability with a decay rate per cycle of 0.054% after 1000 cycles at 1C and outstanding rate performance (703.3 mAh g⁻¹ at 3C). Furthermore, it achieves an impressive areal capacity of 5.64 mAh cm⁻² at 0.2C even with a high sulfur loading of 5.5 mg cm⁻², demonstrating its potential for practical, high-energy-density applications in lithium–sulfur batteries.

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1. Introduction

Lithium–sulfur batteries are considered as the next-generation energy storage system owing to their excellent theoretical specific capacity (1675 mAh g⁻¹) and energy density (2600 Wh kg⁻¹).^1–3 However, the poor conductivity of sulfur and its discharge products (Li₂S/Li₂S₂) results in a decline in sulfur utilization, and the volume expansion in the sulfur cathode also undermines the stability of the electrode structure.^4–6 Besides, the shuttling effect and slow redox kinetic process of Li polysulfides (LiPSs) present substantial barriers to their application in lithium–sulfur batteries.^7–9 Therefore, introducing host materials to modify sulfur cathodes is an efficient strategy to realize further development in lithium–sulfur batteries.^10,11

Researchers have made many attempts to design and synthesize advanced sulfur host materials to solve the above problems in the sulfur cathode.^12–15 Carbon materials, including carbon nanotubes, graphene and conductive carbon black, are widely used as sulfur host materials^16,17 because of their high electronic conductivity, good compatibility with sulfur and excellent chemical stability.^18–20 Nevertheless, non-polar carbon materials find it difficult to inhibit the shuttle effect effectively because of their weak interactions with LiPSs, ultimately leading to severe capacity decay.^21–23 The strategic doping modification of carbon-based materials is a useful approach for improving the adsorption capacity. For example, the incorporation of non-metallic heteroatoms, such as N and O, into the carbon matrix not only enhances its intrinsic electronic conductivity but also significantly increases its polarity, greatly improving its adsorption capacity and mitigating the shuttle effect.^24–27 Furthermore, researchers have frequently integrated polar metal compounds with carbon-based frameworks to obtain advanced composite materials, aiming to synergistically enhance the chemical adsorption and catalytic ability of LiPSs.^28–30 This strategic modification improves the conductivity of carbon while introducing strong polar interactions that effectively adsorb and inhibit LiPSs. Simultaneously, the catalytic activity of these metal compounds accelerates the redox kinetics of LiPS conversion, which mitigates the sluggish reaction rates and suppresses the shuttle effect, thereby optimizing the overall electrochemical performance of lithium–sulfur batteries.^31–34

As a carbon precursor, pollen-derived carbon has the following advantages: (1) intrinsic elemental composition of pollen: owing to its richness in various non-metallic elements, PC undergoes direct in situ modification during carbonization, which produces a conductive and polar carbon matrix without additional doping steps. (2) Hierarchical porosity and structural stability: the naturally self-organized porous network of pollen-derived carbon facilitates efficient sulfur loading while simultaneously alleviating the volumetric expansion of sulfur during lithiation and maintaining structural stability. (3) Sustainable and green chemistry approach: as an abundant and renewable biomass material, PC avoids special and complex synthesis routes, reflecting the advantages of green chemistry.^35,36 In order to further improve the adsorption and catalytic performance of carbon materials, the incorporation of metallic compounds is often necessary to form an active interface. Among the available materials, cerium dioxide (CeO₂) has received significant attention due to its extraordinary properties:^14,37,38 (1) CeO₂ possesses a unique Ce³⁺/Ce⁴⁺ redox couple, which actively participates in accelerating the reversible catalytic conversion of LiPSs. (2) Abundant oxygen vacancies in CeO₂ serve as active sites to enhance the chemisorption and catalytic conversion of LiPSs.

In this study, to enhance catalytic efficiency, abundant electron-delocalized CeO₂ nanoparticles anchored on the surface of pollen-derived carbon materials (PC-CeO₂) are prepared through a carbonization process. The obtained PC-CeO₂ reduced the volume expansion, which maintained the stability and integrity of the structure due to its special and robust structure. Meanwhile, it provided sufficient active sites of electron delocalization for the chemical adsorption and catalytic kinetic conversion of LiPSs. DFT theoretical calculations further prove the strong adsorption and catalytic capacity of PC-CeO₂ towards sulfur species. As a result, the PC-CeO₂/S cathode demonstrated outstanding electrochemical performance with a low fading rate of 0.054% per cycle at 1C after 1000 cycles and an outstanding rate capacity of 703.3 mAh g⁻¹ at 3C. Even with a high sulfur loading of 5.5 mg cm⁻², the PC-CeO₂/S electrode exhibited a high capacity of 5.64 mAh cm⁻² at 0.2C after 60 cycles, which reached commercial standards of 4.0 mAh cm⁻².

2. Results and discussion

The synthesis and preparation process of pollen-derived porous carbon anchored with CeO₂ NPs are illustrated in Fig. 1a. Initially, the pollen is treated with water and acetone to remove fats, sugars and other substances. Then, the as-synthesized CeO₂ NPs are anchored on the surface of the pollen-derived carbon upon high-temperature carbonization at 900 °C to achieve PC-CeO₂. The PC-CeO₂/S is further obtained after the sulfur loading treatment. As shown in the TEM image (Fig. 1b), the CeO₂ NPs display a highly uniform size distribution across the samples, which are uniformly distributed on the PC surface with a porous skeleton (Fig. 1c and S1a–c), The HRTEM image of PC-CeO₂ (Fig. 1d) distinctly reveals a lattice spacing of 0.31 nm, which corresponds to the (111) plane of CeO₂, verifying the structural integrity of the CeO₂ NPs within the composite.³⁹ Furthermore, the EDS elemental mapping images (Fig. 1e) demonstrate the uniform presence of Ce, O, C and N elements within PC-CeO₂. Besides, as shown in Fig. S2a and b, both the interior and surface of PC-CeO₂ and PC are filled with sulfur after sulfur fixation. In addition, compared with PC, the incorporation of CeO₂ NPs into PC results in a reduction in the specific surface area (9.56 m² g⁻¹ for PC-CeO₂ vs. 20.60 m² g⁻¹ for PC, as shown in Fig. S3a and b), which leads to the formation of a hybrid holey structure with an amount of mesopores (Fig. S4a and b). The above results verify the successful integration and structural modification of PC with CeO₂ NPs. After the sulfur is loaded, the specific surface areas of both PC-CeO₂ and PC decrease, as shown in Fig. S5a and b. Furthermore, PC-CeO₂/S exhibits obvious mesoporous pores in Fig. S5c, which contributes to the inhibition of the shuttle effect of LiPSs.

Fig. 1 (a) The preparation process of PC-CeO₂ and PC-CeO₂/S; (b) TEM image of the CeO₂ nanoparticles; (c) SEM image of PC-CeO₂; (d) HRTEM image of PC-CeO₂; (e) the corresponding elemental mapping images of Ce, O, C and N in PC-CeO₂; (f) XRD patterns of PC, CeO₂ and PC-CeO₂; (g) Raman spectra of PC and PC-CeO₂; (h) Ce 3d and (i) O 1s XPS spectra of PC-CeO₂; and (j) EPR curve of PC-CeO₂.

To further confirm the composition and structural characteristics of the material, the relevant characterization results are displayed in Fig. 1f–j. In the XRD patterns of PC-CeO₂ and CeO₂ (Fig. 1f), the diffraction peaks at 28.5°, 33.1°, 47.5° and 56.3° are attributed to the (111), (200), (220) and (311) planes of CeO₂ (PDF#34-0394), respectively, while the XRD pattern of PC shows two distinct carbon peaks.⁴⁰ To investigate the configuration of carbon in PC-CeO₂ and PC, the Raman spectrum was carried out, as shown in Fig. 1g. Two prominent peaks emerge at 1350 and 1580 cm⁻¹, corresponding to the D and G bands of the carbon layer, respectively. Besides, the intensity ratios (I_D/I_G) for PC-CeO₂ and PC are estimated to be 0.98 and 0.96, respectively. This marked difference underscores more defects exhibited by the PC-CeO₂ composite, which plays a pivotal role in optimizing the rapid transport of electron ions during electrochemical process, thereby improving the electrochemical performance. In addition, XPS measurements were carried out to determine the surface chemistry phase and composition of the PC-CeO₂ composite. The XPS full spectrum of PC-CeO₂ verifies the presence of Ce, O, N and C elements (Fig. S6a). As depicted in Fig. S6b, the C 1s spectrum displays four distinct peaks at 284.8 eV (C–C), 286.1 eV (C–N), 287.2 eV (C–O) and 289.2 eV (C=O). Fig. S6c shows three peaks, including pyridinic-N (398.4 eV), pyrrolic-N (400.7 eV) and graphitic-N (401.7 eV). As shown in Fig. 1h, the Ce 3d spectrum exhibits eight distinct splitting peaks. The peaks at 917.1, 908.2, 901.2, 898.6, 888.6, and 882.7 eV can be ascribed to Ce⁴⁺, while others located at 904.1 and 885.7 eV belong to Ce³⁺. In addition, the oxygen vacancy of PC-CeO₂ can be verified by O1s XPS (Fig. 1i). The peaks at 529.5 eV are connected to metal oxygen, and the peaks at 531.5 eV are attributed to a low-coordinated oxygen vacancy.⁴¹ EPR is used to further estimate the oxygen vacancy of PC-CeO₂ (Fig. 1j). The representative EPR signal at g = 2.002 reveals electron capture at the oxygen vacancy.^42,43 The above results explain that the addition of CeO₂ NPs plays a pivotal role in promoting chemical adsorption and accelerating the catalytic conversion of LiPSs.^42,44 The sulfur contents in PC-CeO₂/S, CeO₂/S and PC/S are 75, 71.2 and 74.2 wt%, respectively (Fig. S7). Besides, the EDS elemental mapping of the PC-CeO₂/S in Fig. S8 confirms the uniform distribution of sulfur. These findings demonstrate the remarkable sulfur-holding ability of the PC-CeO₂ composite.

To further demonstrate the adsorption capacity of different materials for LiPSs, visual adsorption experiments were carried out. As illustrated in Fig. 2a, the color in Li₂S₆ with PC-CeO₂ nearly vanishes after 2 h, and the UV-vis data further confirm its exceptional adsorption ability. Furthermore, XPS analysis of LiPS adsorption was performed to analyze the possible mechanism of the strong anchoring effect for LiPSs on PC-CeO₂, as shown in Fig. 2b and c. After exposure to Li₂S₆, the Ce 3d peaks exhibit an obvious shift toward lower binding energies, indicating that Ce cations actively participate in the redox interactions. This shift confirms the involvement of the Ce⁴⁺/Ce³⁺ redox couple, highlighting their crucial role as electron transfer mediators during the Li₂S₆ adsorption process. As depicted in Fig. 2c, the peaks at 165.2 and 164.0 eV correspond to S₂²⁻, and the signals at 168.4 and 169.4 eV correspond to SO₃²⁻ and SO₄²⁻ (thiosulfate), respectively.⁴⁴ All the above data indicate a strong interaction between PC-CeO₂ and Li₂S₆. Li₂S plays a crucial role in the reaction process of lithium–sulfur batteries, affecting the charge and discharge capacity of the batteries. Thus, nucleation (Fig. 2d–f) and dissolution experiments of Li₂S (Fig. 2g–i) were carried out. The results indicate that the PC-CeO₂ electrode exhibits the earliest current peak at 2130 s and the largest deposition capacity at 189.98 mAh g⁻¹. Besides, as shown in Fig. 2g–i, the dissolution profiles of Li₂S display that the PC-CeO₂ electrode exhibits the earliest current response and largest current peak (829.9 s/0.498 mA). Therefore, PC-CeO₂ significantly enhances liquid–solid reactions and accelerates the conversion of LiPSs in Li–S batteries.

Fig. 2 (a) Visual adsorption tests of pure Li₂S₆ solutions, PC + Li₂S₆, CeO₂ + Li₂S₆ and PC-CeO₂ + Li₂S₆. (b) Ce 3d XPS spectra of PC-CeO₂ and PC-CeO₂ + Li₂S₆. (c) S 2p XPS spectra of PC-CeO₂ + Li₂S₆. Precipitation profiles of Li₂S with (d) PC-CeO₂, (e) CeO₂ and (f) PC. Dissolution profiles of Li₂S with (g) PC-CeO₂, (h) CeO₂ and (i) PC.

The cyclic voltammetry (CV) profiles for Li–S batteries equipped with PC-CeO2, CeO2 and PC electrodes are depicted in Fig. 3a. Two notable reduction peaks, labeled peaks A and B, appear around 2.30 V and 2.02 V, respectively. These correspond to the stepwise reduction process of S8 to lithium polysulfides (Li2Sn, 4 ≤ n ≤ 8) and their subsequent transformation into shorter, insoluble species, like Li2S2 and Li2S. The oxidation peak, denoted as peak C at 2.31 V, marks the reverse reaction, where Li2S2 and Li2S are oxidized back to S8. Among these, the PC-CeO2 electrode stands out with the highest current response and the smallest reaction polarization (0.295 V for PC-CeO2 vs. 0.314 V for CeO2 and 0.358 V for PC), signifying its superior electrochemical performance in Li–S batteries. Fig. 3b and c illustrate the enlarged CV curve segments around peaks A and C in Fig. 3a. Tafel plots were derived based on these peaks, revealing that the PC-CeO2 electrode exhibits lower Tafel slopes (27.3 mV dec−1 for peak A and 98.91 mV dec−1 for peak C) than the CeO2 electrode (28.2 mV dec−1 and 124.53 mV dec−1) and PC electrode (54.9 mV dec−1 and 111.38 mV dec−1). This suggests the fastest electron transfer rate of PC-CeO2 during the redox process of LiPSs. To verify the catalytic capabilities of the different electrodes in facilitating LiPS conversion, symmetric cells were assembled. Fig. 3d shows that the PC-CeO2 electrode demonstrates the highest current density and the largest curve area, underscoring its excellent catalytic activity in the conversion of LiPSs. In addition, galvanostatic intermittent titration technique (GITT) profiles were measured at 0.1C to explore the kinetic performance of the host materials (Fig. 3e and S9a–f). Two important transition points about Li2S nucleation and activation are shown in Fig. 3e, and the voltage difference at Li2S nucleation and Li2S activation points can be calculated (Fig. 3f). The cell with the PC-CeO2 electrode shows the lowest Li2S nucleation voltage difference (22.8 mV) and Li2S activation voltage difference (79.1 mV), indicating its lower polarization and faster kinetics. An EIS test was conducted to analyze the redox kinetics (Fig. 3g). The resulting Nyquist plots are presented in Fig. 3g. Notably, the PC-CeO2 cathode exhibits a relatively low interface resistance (Rs) of 2.72 Ω, compared to 4.74 Ω for the CeO2 cathode and 2.78 Ω for the PC cathode. Additionally, the PC-CeO2 cathode displays a relatively small charge transfer resistance (Rct) of 63.9 Ω, in contrast to 81.5 Ω for the CeO2 and a significantly high value of 167.5 Ω for the PC cathode (shown in Table S2). These results indicate that PC-CeO2 facilitates the rapid kinetics conversion of LiPSs in Li–S batteries. Furthermore, the CV curves of Li–S cells with different cathodes were recorded over a scan range of 0.1–0.5 mV s−1 (Fig. 3h, S10a and S10b). The CV tests were conducted 15 times from 0.1 to 0.5 mV s−1, and the CV curves at the same scanning speed are very close, as depicted in Fig. S11, demonstrating the exceptional electrochemical stability and reversibility of the PC-CeO2 electrode. The peak current is plotted against the square root of the scan rate, as shown in Fig. 4i. According to the classical Randles–Sevcik equation (IP = 2.69 × 105n1.5ADLi+0.5CLi+ν0.5), the Li+ diffusion coefficients are calculated. For the PC-CeO2 electrode, the Li+ diffusion rates are found to be DLi+ (A) = 2.35 × 10−8 cm2 s−1, DLi+ (B) = 4.9 × 10−8 cm2 s−1 and DLi+ (C) = 5.21 × 10−8 cm2 s−1, which are higher than those of the CeO2 and PC electrodes (Fig. S12). These results demonstrate that the PC-CeO2 electrode delivers the highest Li+ diffusion coefficients, leading to faster electrochemical reactions and enhanced performance in Li–S batteries.

Fig. 3 Electrochemical performances of the PC-CeO₂, CeO₂ and PC electrodes: (a) CV curves at 0.1 mV s⁻¹. (b and c) The relevant section of the CV curves (peaks A and C) and the Tafel plots (Insets). (d) CV curves of the symmetric cells with different electrodes. (e) GITT curve of the PC-CeO₂ electrode. (f) The potential difference at Li₂S nucleation and Li₂S activation point. (g) EIS spectrum. (h) CV profiles of the PC-CeO₂ electrode from 0.1 to 0.5 mV s⁻¹. (i) The corresponding linear fitting of the peak current.

Fig. 4 Electrochemical performances of the PC-CeO₂, CeO₂ and PC electrodes: (a) charging/discharging profiles for the initial cycle at 0.2C. (b) ΔE and Q₂/Q₁ values obtained from discharge/charge curves. (c) rate capability at 0.2–3 C. (d) cycling stability of the PC-CeO₂ electrode at 0.2C with a sulfur loading of 5.5 mg cm⁻². (e) Cycling performance at 1C. (f) Cycling performance of the pouch cell assembled with the PC-CeO₂ electrode at 0.2C. (g) Electrochemical performances of different biomass materials.

Li–S batteries were assembled using PC-CeO₂, CeO₂ and PC electrodes paired with Li anodes. The electrochemical results are summarized in Fig. 4. The charge/discharge profiles show two distinct discharge plateaus and one charge plateau at 0.2C (Fig. 4a), consistent with the CV curve data. Besides, the PC-CeO₂ electrode exhibits a smaller polarization potential (ΔE = 182 mV) than the CeO₂ electrode (ΔE = 278 mV) and PC electrode (ΔE = 207 mV), revealing its enhanced redox kinetics and reversibility in Li–S batteries (Fig. 4b). Moreover, the capacity ratio between the two discharge plateaus (Q₂ and Q₁) serves as a crucial indicator of catalytic activity. Q₁ corresponds to sulfur reduction to soluble LiPSs, while Q₂ reflects their conversion into Li₂S. The PC-CeO₂ electrode shows the highest Q₂/Q₁ ratio of 2.17 (vs. 2.07 for CeO₂ and 2.15 for PC), suggesting more efficient sulfur utilization and superior catalytic activity. The cycling performance of the electrodes at 0.2C is depicted in Fig. S13, where the PC-CeO₂ cathode displays an initial discharge capacity of 1174 mAh g⁻¹, with a coulombic efficiency (CE) consistently above 99.8% over 200 cycles, surpassing both the CeO₂ and PC cathodes.

As shown in Fig. 4c, the rate capabilities were further evaluated from 0.2 to 3C. The PC-CeO₂ electrode demonstrates a remarkable rate performance, maintaining initial discharge capacities of 1024.4, 891.5, 829.1, 744.8, and 703.3 mAh g⁻¹ at 0.2, 0.5, 1, 2, and 3C, respectively. Even when the current density is reversed back to lower rates (2, 1, 0.5, and 0.2C), the discharge capacities remain at 756.8, 827.3, 883.2, and 957.8 mAh g⁻¹, achieving capacity retention rates of 101.6% at 2C, 99.7% at 1C, 99.1% at 0.5C, and 93.5% at 0.2C, respectively. The corresponding galvanostatic discharge/charge voltage curves of PC-CeO₂, CeO₂ and PC at 0.2–3 C are presented in Fig. S14a–c. The PC-CeO₂ electrode maintains a stable discharge/charge voltage plateau at 3C owing to its enhanced ionic/electronic conductivity, which accelerates the redox reactions of LiPSs. Besides, the PC-CeO₂ cathode delivers an impressive initial specific areal capacity of 7.5 mAh cm⁻² and maintains a reversible capacity of 5.6 mAh cm⁻² after 60 cycles at 0.2C at a sulfur loading of 5.5 mg cm⁻² (Fig. 4d). The long-term cycling stability was tested at 1C (Fig. 4g), where the PC-CeO₂ cathode achieved an initial discharge capacity of 826.2 mAh g⁻¹, gradually declining to 375.1 mAh g⁻¹ after 1000 cycles, with an exceptionally low capacity decay rate of just 0.054% per cycle and 0.035% per cycle at 2C, as depicted in Fig. S15. This demonstrates the remarkable cycling stability of the PC-CeO₂ cathode. A pouch cell was constructed to assess the practical viability of the PC-CeO₂ electrode. As displayed in Fig. 4f, the battery exhibits an impressive initial specific capacity of 845.2 mAh g⁻¹ at 0.2C. Furthermore, as demonstrated in Fig. S16, a rigorous test of the pouch battery's mechanical properties was conducted by folding it at various angles. Despite these deformations, the battery effortlessly illuminates a “NENU” panel, showing its remarkable anti-folding characteristics and mechanical durability. This underscores the potential of Li–S batteries with PC-CeO₂ cathode for flexible and reliable energy storage applications. In order to further verify the influence of different host materials in Li–S batteries, the cells with the PC-CeO₂ and PC electrodes are disassembled after 500 cycles at 2C, and the SEM images are shown in Fig. S17. First, the SEM images of the sulfur cathodes paired with the PC-CeO₂ and PC electrodes are shown in Fig. S17a and b, respectively. Compared with the cathode assembled with the PC-CeO₂ electrode, the cathode with the PC electrode is severely split after 500 cycles. Besides, the SEM images of Li anodes show that the lithium electrode assembled with the PC-CeO₂ electrode is well protected with a smooth surface and uniform deposition (Fig. S17c). On the contrary, the surface of the lithium electrode with the PC electrode is severely damaged after 500 cycles (Fig. S17d). As illustrated in Fig. S18, the EDS mapping on the cycled separator (a) and cathode (b) shows the uniform distribution of Ce, S, O and N elements. Apart from that, in the XRD patterns of PC-CeO₂ cathode (Fig. S19), the diffraction peak at 27° is indexed to the (111) planes of Li₂S (PDF#26-1188), which emphasizes the accumulation of Li₂S. In conclusion, the cells with the PC-CeO₂ host material can not only inhibit the shuttle of LiPSs but also adsorb LiPSs and catalyze the conversion. Meanwhile, as illustrated in Fig. 4g and Tables S3 and 4, the electrochemical performance of PC-CeO₂ is more competitive than other sulfur host materials in Li–S batteries.^45–50

To further explain the adsorption and catalytic capacity of LiPSs for the PC-CeO₂, the density functional theoretical calculations were conducted, and the (111) plane and (110) planes of CeO₂ were chosen as the calculation plane owing to the HRTEM image and the XRD results of PC-CeO₂. The corresponding results are presented in Fig. 5a–c. The adsorption models of sulfur species on PC-CeO₂(111), PC-CeO₂(110) and PC are shown in Fig. 5a, S20 and S21, respectively. As shown in Fig. 5b, PC-CeO₂(111) exhibits the highest adsorption energy towards different sulfur species compared with the PC, and the adsorption energies among Li₂S, Li₂S₂, Li₂S₄, Li₂S₆, Li₂S₈, S₈ and PC-CeO₂ are calculated as −4.61, −3.73, −2.71, −2.48, −2.73 and −1.10 eV, respectively. The above results further illustrate the strong adsorption capacity between PC-CeO₂(111) and LiPSs. Fig. 5c depicts the Gibbs free energy curves for the reduction of sulfur from S₈ to Li₂S on PC-CeO₂(111), PC-CeO₂(110) and PC. The process from S₈ to Li₂S₈ for PC-CeO₂(111), PC-CeO₂(110) and PC is calculated as −4.47, −3.68 and −2.82 eV, respectively, showing that the step is a spontaneous reaction, and the results further show that the process from S₈ to Li₂S₈ for PC-CeO₂(111) is the easiest to proceed. Besides, the highest positive Gibbs free energy (ΔG) in the entire sulfur reduction process is the conversion from Li₂S₂ to Li₂S, marking it the rate-limiting step in the LiPS reduction reaction. Notably, the ΔG for PC-CeO₂(111) is only 0.05 eV, which is lower than 0.62 eV for PC-CeO₂(110) and 1.11 eV for PC, indicating that the PC-CeO₂(111) can accelerate the electrochemical redox reactions of sulfur. Finally, the mechanism of the PC-CeO₂/S cathode to enhance reaction kinetics in Li–S batteries is illustrated in Fig. 5d. On the one hand, the 3D porous structure offers ample space for the loading of CeO₂ NPs as active sites, further offering enough chemisorption sites for LiPSs. On the other hand, the CeO₂ shows remarkable adsorption and catalysis ability to LiPSs, improving the cycle stability and electrochemical performance of the batteries significantly.

Fig. 5 (a) The corresponding model models of sulfur and LiPSs on the CeO₂ (111) plane. (b) The adsorption energies of sulfur species on the PC-CeO₂ (111), PC-CeO₂ (110) and PC planes. (c) The Gibbs free energies from S₈ to Li₂S on the PC-CeO₂ (111), PC-CeO₂ (110) and PC planes. (d) The mechanism of the PC-CeO₂/S cathode to promote the LiPS conversion.

3. Conclusion

In summary, the electron-delocalized CeO₂ nanoparticles were anchored on the surface of pollen-derived carbon and acted as host materials for lithium–sulfur batteries. Specifically, the PC-CeO₂ as the sulfur host highlights the following advantages. First, the pollen-derived porous carbon skeleton as the main body of the host accelerates the transfer of electrons and ions, alleviates volume expansion and maintains the structural stability of the sulfur cathode. Second, the introduction and modification of CeO₂ increase the electron delocalization to enhance the chemical adsorption of LiPSs, effectively inhibiting the shuttle of LiPSs and catalyzing its reaction process. Third, owing to the intrinsic structure and properties of these two components, PC-CeO₂ can foster a profound synergistic interaction while maintaining the stability of the overall structure. Finally, the PC-CeO₂/S cathode displays exceptional electrochemical performance, which exhibits a capacity decay rate of 0.054% at 1C after 1000 cycles and an outstanding rate performance of 703.3 mAh g⁻¹at 3C, as well as an excellent high capacity of 5.64 mAh cm⁻² at 0.2C. This study provides a feasible strategy for the development of biomass and its composites as sulfur hosts in lithium–sulfur batteries.

Conflicts of interest

The authors declare no conflict of interest.

Data availability

The data supporting this article are included in the supplementary information (SI). Supplementary information: experimental section, additional Fig. S1–S21 and Tables S1–4 can be found in supporting information.). See DOI: https://doi.org/10.1039/d5ta08719h.

Acknowledgements

This work is financially supported by the Jilin Provincial Science and Technology Development Foundation (Grant No. 20250102063JC), the Natural Science Foundation of Jiangsu Province (BK 20210130), and the Opening funding from Key Laboratory of Engineering Dielectrics and Its Application (Harbin University of Science and Technology) (No. KFM202507, Ministry of Education) as well as Open Project of Inner Mongolia Engineering Research Center of Lithium–Sulfur Battery Energy Storage (No. MDK2025071). Jian Wang acknowledges the funding provided by the Alexander von Humboldt Foundation. The High-Performance Computing Center of Xiamen University is acknowledged for its supercomputer resources.

References

1. X. Li, Q. Guan, Z. Zhuang, Y. Zhang, Y. Lin, J. Wang, C. Shen, H. Lin, Y. Wang, L. Zhan and L. Ling, ACS Nano, 2023, 17, 1653–1662.
2. J. Wang, J. Zhang, Y. Zhang, H. Li, P. Chen, C. You, M. Liu, H. Lin and S. Passerini, Adv. Mater., 2024, 36, 2402792.
3. J. Wang, J. Zhang, J. Wu, M. Huang, L. Jia, L. Li, Y. Zhang, H. Hu, F. Liu, Q. Guan, M. Liu, H. Adenusi, H. Lin and S. Passerini, Adv. Mater., 2023, 35, e2302828.
4. H. Ma, Z. Yu, H. Li, D. Guo, Z. Zhou, H. Jin, L. Wu, X. a. Chen and S. Wang, Adv. Funct. Mater., 2024, 34, 2310301.
5. F. Zhou, Y. Meng, T. Wang, D. Sun, L. Gao, Z. Sun, Y. Wang, J. Zeng, B. Wang, R. Zhang, R. Li, C. Huang, L. Yang and X. Wang, Nano Energy, 2024, 127, 109755.
6. J. Zhang, C. You, H. Lin and J. Wang, Energy Environ. Mater., 2022, 5, 731–750.
7. L. Li, H. Tu, J. Wang, M. Wang, W. Li, X. Li, F. Ye, Q. Guan, F. Zhu, Y. Zhang, Y. Hu, C. Yan, H. Lin and M. Liu, Adv. Funct. Mater., 2023, 33, 2212499.
8. X. Zhang, X. Li, Y. Zhang, X. Li, Q. Guan, J. Wang, Z. Zhuang, Q. Zhuang, X. Cheng, H. Liu, J. Zhang, C. Shen, H. Lin, Y. Wang, L. Zhan and L. Ling, Adv. Funct. Mater., 2023, 33, 2302624.
9. J. Wang, J. Zhang, S. Duan, L. Jia, Q. Xiao, H. Liu, H. Hu, S. Cheng, Z. Zhang, L. Li, W. Duan, Y. Zhang and H. Lin, Nano Lett., 2022, 22, 8008–8017.
10. Z. Yang, Z. Hu, G. Yan, M. Li, Y. Feng, X. Qu and X. Zhang, J. Colloid Interface Sci., 2023, 629, 65–75.
11. L. Yoshida, T. Hakari, Y. Matsui, M. Deguchi, H. Yamamoto, M. Inoue and M. Ishikawa, J. Power Sources, 2024, 624, 235572.
12. M. Wang, Z. Bai, T. Yang, C. Nie, X. Xu, Y. Wang, J. Yang, S. Dou and N. Wang, Adv. Energy Mater., 2022, 12, 2201585.
13. H. Li, H. Yang and X. Ai, Adv. Mater., 2023, 2305038.
14. J. Zhang, R. He, Q. Zhuang, X. Ma, C. You, Q. Hao, L. Li, S. Cheng, L. Lei, B. Deng, X. Li, H. Lin and J. Wang, Adv. Sci., 2022, 9, e2202244.
15. J. Wang, L. Jia, H. Liu, C. Wang, J. Zhong, Q. Xiao, J. Yang, S. Duan, K. Feng, N. Liu, W. Duan, H. Lin and Y. Zhang, ACS Appl. Mater. Interfaces, 2020, 12, 12727–12735.
16. L. Jia, J. Wang, Z. Chen, Y. Su, W. Zhao, D. Wang, Y. Wei, K. Jiang, J. Wang, Y. Wu, J. Li, W. Duan, S. Fan and Y. Zhang, Nano Res., 2019, 12, 1105–1113.
17. J. Wang, J. Zhang, S. Duan, T. Li, L. Jia, H. Liu, L. Li, S. Cheng, H. Hu, M. Huang, H. Hu, S. Zhang, Q. Xiao and H. Lin, Chem. Eng. J., 2022, 429, 132352.
18. T. Yang, J. Xia, Z. Piao, L. Yang, S. Zhang, Y. Xing and G. Zhou, ACS Nano, 2021, 15, 13901–13923.
19. D.-S. Kim, S.-G. Woo, C.-J. Kang, J.-H. Lee, J.-N. Lee, J.-S. Yu and Y.-J. Kim, ChemSusChem, 2023, 16, e202202009.
20. M. Zheng, Y. Chi, Q. Hu, H. Tang, X. Jiang, L. Zhang, S. Zhang, H. Pang and Q. Xu, J. Mater. Chem. A, 2019, 7, 17204–17241.
21. B.-J. Lee, T.-H. Kang, H.-Y. Lee, J. S. Samdani, Y. Jung, C. Zhang, Z. Yu, G.-L. Xu, L. Cheng, S. Byun, Y. M. Lee, K. Amine and J.-S. Yu, Adv. Energy Mater., 2020, 10, 1903934.
22. H. Wang, W. Zhang, J. Xu and Z. Guo, Adv. Funct. Mater., 2018, 28, 1707520.
23. J. Wang, L. Jia, J. Zhong, Q. Xiao, C. Wang, K. Zang, H. Liu, H. Zheng, J. Luo, J. Yang, H. Fan, W. Duan, Y. Wu, H. Lin and Y. Zhang, Energy Storage Mater., 2019, 18, 246–252.
24. L. Wang, X. Meng, X. Wang and M. Zhen, Small, 2023, 19, 2300089.
25. P. Chen, T. Wang, D. He, T. Shi, M. Chen, K. Fang, H. Lin, J. Wang, C. Wang and H. Pang, Angew. Chem., Int. Ed., 2023, 62, e202311693.
26. J. Zhang, C. You, H. Lin and J. Wang, Energy Ennviron. Mater, 2022, 5, 731–750.
27. J. Wang, J. Zhang, S. Cheng, J. Yang, Y. Xi, X. Hou, Q. Xiao and H. Lin, Nano Lett., 2021, 21, 3245–3253.
28. X. Ren, Q. Wang, Y. Pu, Q. Sun, W. Sun and L. Lu, Adv. Mater., 2023, 35, 2304120.
29. J. Cho, S. Ryu, Y. J. Gong, S. Pyo, H. Yun, H. Kim, J. Lee, J. Yoo and Y. S. Kim, Chem. Eng. J., 2022, 439, 135568.
30. K. Zou, W. Jing, X. Dai, X. Chen, M. Shi, Z. Yao, T. Zhu, J. Sun, Y. Chen, Y. Liu and Y. Liu, Small, 2022, 18, 2107380.
31. X. Wang, G. Li, M. Li, R. Liu, H. Li, T. Li, M. Sun, Y. Deng, M. Feng and Z. Chen, J. Energy Chemistry, 2021, 53, 234–240.
32. J. Luo, K. Guan, W. Lei, S. Zhang, Q. Jia and H. Zhang, J. Energy Chem., 2022, 122, 101–120.
33. Z. Yang, X. Liang, K. Chen, G. Zhang and S. Xing, J. Energy Storage, 2024, 102, 114218.
34. J. Wang, L. Jia, S. Duan, H. Liu, Q. Xiao, T. Li, H. Fan, K. Feng, J. Yang, Q. Wang, M. Liu, J. Zhong, W. Duan, H. Lin and Y. Zhang, Energy Storage Mater., 2020, 28, 375–382.
35. H. Hu, L. Cui, W. Gao, Z. Zhang and S. Kang, Chem. Eng. J., 2023, 453, 139516.
36. Q. Xiao, G. Li, M. Li, R. Liu, H. Li, P. Ren, Y. Dong, M. Feng and Z. Chen, J. Energy Chem., 2020, 44, 61–67.
37. G. Wen, X. Zhang, Z. Shi, Y. Sui, J. Li, J. Zeng, J. Zheng and L. Wu, J. Colloid Interface Sci., 2022, 619, 106–115.
38. J. Dong, X. Cheng, H. Yang, H. Li, H. Liu, L. Jia, Y. Zhang, Q. Guan, J. Jia, F. Wu, J. Zhang, M. Liu, H. Lin and J. Wang, Adv. Mater., 2025, e2501079,  DOI:10.1002/adma.202501079.
39. J. Zhang, Q. Rao, B. Jin, J. Lu, Q.-g. He, Y. Hou, Z. Li, X. Zhan, F. Chen and Q. Zhang, Chem. Eng. J., 2020, 388, 124120.
40. G. Wen, Z. Shi, Y. Sui, B. Wang, X. Zhang, Z. Zhang and L. Wu, J. Colloid Interface Sci., 2022, 623, 697–702.
41. N. P. d. A. Mendes, A. L. d. Souto Neto, J. d. S. Hortêncio, A. L. Menezes de Oliveira, R. A. Raimundo, D. A. Macedo and F. F. da Silva, ACS Omega, 2024, 9, 49913–49924.
42. K. Zou, X. Chen, W. Jing, X. Dai, P. Wang, Y. Liu, R. Qiao, M. Shi, Y. Chen, J. Sun and Y. Liu, Energy Storage Mater., 2022, 48, 133–144.
43. Y. Song, H. Li, J. Li, J. An, J.-J. Shao and G. Zhou, J. Energy Chem., 2023, 87, 51–60.
44. H. Jia, J. Fan, P. Su, T. Guo and M.-C. Liu, Small, 2024, 20, 2311343.
45. I. Marri, Z. Edfouf, A. Caballero and A. Y. Tesio, J. Energy Storage, 2024, 82, 110518.
46. X. Wang, M. Niu, C. Gao, Y. Li, T. Li and B. Ren, Carbon Lett., 2024, 34, 1385–1398.
47. J. Xu, P. Zhou, L. Dai, Y. Gui, L. Yuan, X. Shen, C. Zhang and K. Huo, Green Chem., 2021, 23, 6008–6019.
48. G. Liu, J. Yang, Y. Wang, H. Lin, M. Zheng, H. Hu, H. Dong, X. Liu, Y. Liu and Y. Xiao, J. Energy Storage, 2023, 67, 107466.
49. F. Liu, P. Feng, M. Yuan, G. Zhai, M. T. Innocent, H. Xiang, Q. Wu, Y. Lu and M. Zhu, ACS Sustain. Chem. Eng., 2023, 11, 16544–16553.
50. S.-Y. Lee, Y. Choi, J.-K. Kim, S.-J. Lee, J. S. Bae, E. D. Jeong and J. Indus, J. Ind. Eng. Chem., 2021, 94, 272–281.

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Footnote

† These authors contributed equally to this work.

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