Ionic‐Liquid‐Triggered Amorphization Engineers Symmetry‐Breaking p‐Block Bismuth Oxides with Electric Dipole Domains for Practical Lithium‐Sulfur Batteries

ABSTRACT

The practical application of lithium‐sulfur batteries is severely hindered by the sluggish sulfur redox kinetics and the notorious lithium polysulfides (LiPSs) shuttle effect. Herein, we report a strategy utilizing an Fe‐based ionic liquid to trigger amorphization, engineering symmetry‐breaking p‐block bismuth oxides on carbon nanofibers (CNFs) with electric dipole domains and asymmetric Fe₁‐O‐Bi electronic bridges (Fe₁⊂A/C‐Bi₂O₃@CNFs). The amorphous phase induces significant electronic delocalization, facilitating substantial orbital overlap and creating electron transport channels for rapid redox of LiPSs. Specifically, the asymmetric Fe₁‐O‐Bi electron bridges lower the p‐band center through 3d‐2p‐6p multi‐orbital coupling, optimizing the chemical adsorption of LiPSs and preventing active site poisoning. The electronic dipole domain functions as an electron/Li⁺ “pump” to enhance charge transfer and Li⁺ diffusion. In addition, the electric dipole domain induces dipole‐dipole interactions, facilitating Li─S bond polarization and cleavage. As a result, the Fe₁⊂A/C‐Bi₂O₃‐based cell achieved a cyclability of 698 mAh g⁻¹ at 1.0 C over 1000 cycles with a degradation rate of 0.026% per cycle, and a high areal capacity of 6.8 mAh cm⁻² under a sulfur loading of 7.4 mg cm⁻². The strategy of constructing an electronic dipole domain through amorphization provides a new direction for the rational design of efficient catalysts for sulfur redox reactions.

The in situ Fe‐based ionic liquid‐triggered amorphization strategy engineers atomically dispersed Fe₁ within symmetry‐breaking p‐block bismuth oxides on the surface of carbon nanofibers, featuring electric dipole domains and asymmetric Fe₁‐O‐Bi electronic bridges (Fe₁⊂A/C‐Bi₂O₃@CNFs). The Fe₁⊂A/C‐Bi₂O₃@CNFs can effectively anchor lithium polysulfides and accelerate sulfur redox kinetics, thereby enhancing the performance of lithium‐sulfur batteries for practical applications.

1. Introduction

The escalating global demand for high‐energy‐density storage systems, driven by the proliferation of electric vehicles, portable electronics, and grid‐scale renewable energy storage, has underscored the limitations of current lithium‐ion batteries. Lithium‐sulfur (Li||S) batteries feature an impressive theoretical specific capacity of 1675 mAh g⁻¹ and an energy density of approximately 2600 Wh kg⁻¹, making them a promising candidate among next‐generation battery technologies [1]. However, the practical commercialization of Li||S batteries faces several inherent challenges. The insulating nature of sulfur (S₈) and the discharge products (Li₂S/Li₂S₂) lead to poor sulfur redox reaction (SRR) kinetics and low active sulfur utilization [2]. The most critical issue is the notorious “shuttle effect” of soluble lithium polysulfides (LiPSs) [3]. LiPSs are highly susceptible to dissolving in the electrolyte, diffusing between the electrodes, and engaging in parasitic reactions with the lithium metal anode, leading to rapid capacity fade, low Coulombic efficiency, and poor cycle life [4].

Extensive research has been dedicated to mitigating the LiPS shuttle effect. A prominent strategy involves the use of transition metals (TMs)‐based catalysts to chemically anchor and electrocatalytically convert LiPSs [5, 6, 7]. The essential role of TMs catalysts in the SRR arises from their unfilled d‐electron orbitals [8]. LiPSs, acting as Lewis bases, can donate lone electron pairs from sulfur atoms to the vacant d‐orbitals of TMs centers, creating strong chemical adsorption [9, 10]. This interaction effectively traps LiPSs and lowers the energy barriers for the SRR. Nevertheless, the excessive binding of S‐TMs can lead to the irreversible passivation of active sites, poisoning catalytic sites and hindering the subsequent conversion of LiPSs [11, 12, 13].

To address this limitation, attention has shifted to p‐block metal‐based catalysts, such as bismuth (Bi) [14, 15], tin (Sn) [16, 17], and antimony (Sb) [18]. The adsorption capacity and catalytic conversion of LiPSs largely depend on the electronic structure of the catalyst. p‐block metals exhibit a unique electronic structure characterized by fully filled d‐orbitals and chemically active p‐orbitals [19]. The interaction between p‐block metal catalysts and LiPSs primarily involves p‐p orbital hybridization with the sulfur atoms [14]. This interaction is generally more moderate than the d‐p coupling of TMs, effectively avoiding the strong chemisorption that leads to active site poisoning. In addition, the optimized interaction is more conducive to facilitating the reversible conversion of LiPSs. Despite these advantages, p‐block metal catalysts often suffer from intrinsic drawbacks, including a limited number of active sites, relatively low electronic conductivity, and sluggish Li⁺ diffusion, which restrict their catalytic efficiency, especially under high‐sulfur‐loading conditions. Typical strategies for developing high‐performance catalysts involve the creation of defects [20, 21], heteroatom doping [22, 23], and the formation of heterointerfaces [24, 25, 26]. These approaches modify the electronic structures, generating favorable catalytic centers that optimize the adsorption strength for LiPSs and reduce the redox energy barrier. Incorporating vacancies into metal oxides is a common strategy for achieving symmetry breaking [27]. Engineering these defects alters the local coordination environment of atoms surrounding the defect sites, optimizing the chemical anchoring of LiPSs by reconstructing the M─S bonding. In addition, the formation of heterojunctions can induce internal electric fields and adjust the band center of the metal constituents, leading to favorable adsorption strength and improved catalytic activity for SRR [28, 29]. These methodologies disrupt local symmetry to create more favorable electronic configurations for the optimized adsorption of LiPSs [30]. Nonetheless, a significant challenge remains in the rational design of these symmetry‐breaking catalysts to achieve an optimal dynamic adsorption‐desorption equilibrium.

In this work, we introduce an in situ strategy that utilizes an Fe‐based ionic liquid to engineer a novel atomically dispersed Fe₁ crystalline/amorphous p‐block Bi₂O₃, featuring an electronic dipole domain and asymmetric Fe₁‐O‐Bi electronic bridges through amorphous engineering (Fe₁⊂A/C‐Bi₂O₃). Integrated with density functional theory (DFT) calculations and advanced in situ characterizations, Fe₁⊂A/C‐Bi₂O₃ effectively modulates adsorption configurations and lowers the energy barrier for the SRR. The amorphization induces significant electron delocalization, resulting in a diffuse electron cloud over the catalyst that exposes more catalytically active sites and enhances the flexibility of electronic responses. In addition, the difference in Fermi levels between the crystalline and amorphous phases results in a dipole domain with a built‐in electric field, effectively accelerating the electron separation and migration of Li⁺. Importantly, the robust dipole–dipole interactions toward LiPSs are realized to generate localized tensile strain fields to destabilize the S─S/Li─S bonds. The multi‐orbital coupling of the asymmetric Fe₁‐O‐Bi electronic bridge weakens the electron donation of Bi atoms, lowering the p‐band center of the Bi 6p and increasing electron occupancy in the Bi‐S anti‐bonding orbitals. Consequently, Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cells exhibit exceptional cycling stability and remarkable rate performance, even under high sulfur loading conditions. This work provides a novel design strategy for symmetry‐breaking p‐block metal‐based catalysts with an electronic dipole domain via amorphization engineering.

2. Results and Discussion

The fabrication of Fe₁⊂A/C‐Bi₂O₃@CNFs is illustrated schematically in Figure 1a. Black liquor lignin (Figure S1), obtained from the pulping and papermaking process, serves as the carbon source and is further transformed into high‐performance carbon nanofibers (CNFs) (Figure S4) via electrospinning (Figures S2 and S3). Furthermore, atomic‐level Fe₁‐incorporated p‐block bismuth oxide nanosheets with an amorphous‐crystalline structure and asymmetric Fe₁‐O‐Bi electronic bridges are in situ grown on the CNFs through an ionic liquid‐mediated hydrothermal reaction (Figure S5). Similar to C‐Bi₂O₃ (Figure S7), Fe₁⊂A/C‐Bi₂O₃ displays a hollow, flower‐like structure composed of ultra‐thin, 2D petal‐like nanosheets (Figure 1b–d). Notably, both C‐Bi₂O₃ and Fe₁⊂A/C‐Bi₂O₃ could be uniformly dispersed on the surface of the CNFs (Figures S6 and S8). In particular, the Fe₁⊂A/C‐Bi₂O₃@CNFs composite exhibited a significantly larger specific surface area of 84.8 m² g⁻¹ compared to C‐Bi₂O₃@CNFs. Additionally, the distinct Type IV hysteresis loop observed in its adsorption‐desorption isotherm provides compelling evidence for a mesoporous structure (Figure S9). Figure S10 compares the XRD patterns of C‐Bi₂O₃@CNFs and Fe₁⊂A/C‐Bi₂O₃@CNFs. For C‐Bi₂O₃@CNFs, the diffraction peaks located at 2θ = 27.9°, 32.3°, 46.4°, and 55.0° can be indexed to the (111), (200), (220), and (311) planes of the cubic phase (PDF#52‐1007), respectively. In contrast, the (111) reflection of Fe₁⊂A/C‐Bi₂O₃@CNFs exhibits pronounced peak broadening accompanied by a marked intensity attenuation, indicative of suppressed long‐range crystallographic order and a reduced coherently scattering domain size.

FIGURE 1.

(a) Schematic illustration of the Fe₁⊂A/C‐Bi₂O₃@CNFs. (b) SEM and (c, d) TEM images of Fe₁⊂A/C‐Bi₂O₃. (e) HRTEM images, corresponding (f) lattice spacing profiles, (g) fast‐Fourier‐transform images, and (h) crystal plane structures of C‐Bi₂O₃. (i) HRTEM images, (j) corresponding fast‐Fourier‐transform images, and (k) geometric phase analysis of Fe₁⊂A/C‐Bi₂O₃. (l) Spherical aberration‐corrected HAADF‐STEM images and (m) corresponding 3D topographic atom images of Fe₁⊂A/C‐Bi₂O₃. AFM images and corresponding height profiles of (n) Fe₁⊂A/C‐Bi₂O₃ and (o) C‐Bi₂O₃. (p) Elemental distribution of Fe₁⊂A/C‐Bi₂O₃.

High‐resolution transmission electron microscopy (HRTEM) images and the corresponding well‐defined fast Fourier transform (FFT) pattern of C‐Bi₂O₃@CNFs indicate lattice spacings of 2.72 and 3.15 Å associated with the (200) and (111) faces of C‐Bi₂O₃, respectively (Figure 1e–h). Interestingly, the HRTEM image and FFT patterns in Figure 1i,j demonstrate that the Fe₁⊂A/C‐Bi₂O₃@CNFs consist of nanoscale amorphous and crystalline domains, forming an amorphous‐crystalline heterostructure with the amorphous nanodomains embedded in the crystalline matrix. The unique amorphous‐crystalline heterointerface in Fe₁⊂A/C‐Bi₂O₃@CNFs facilitates efficient electronic symmetry breaking, promoting electron delocalization in Bi₂O₃. This, in turn, modulates the p‐band center of the Bi 6p and optimizes the chemical adsorption of LiPSs. In addition, the strain field was derived from the HRTEM image of Fe₁⊂A/C‐Bi₂O₃ using geometric phase analysis, as shown in Figure 1k. The color scale represents the variation in strain, with values ranging from −0.1 to 0.1, as indicated by the color scale bar. Lattice strain shows significant inhomogeneity, with alternating tensile and compressive strain fields. This confirms the presence of microstrain within Fe₁⊂A/C‐Bi₂O₃@CNFs, resulting from atomic size mismatch. Notably, lattice strain can simultaneously modulate the atomic coordination environment of the central metal sites and induce charge rearrangement. Fe‐based ionic liquids can induce a Bi─O mismatch, which distorts the natural (BiO₆) octahedra, disrupting microstructural symmetry. This, in turn, leads to crystal field splitting, achieving a symmetry‐breaking electronic structure in Fe₁⊂A/C‐Bi₂O₃@CNFs. Aberration‐corrected scanning transmission electron microscopy (HAADF‐STEM) was used to visualize the atomic‐level Fe₁ within A/C‐Bi₂O₃. The atomic lattice of Bi atomic columns displayed a uniform distribution, as illustrated in the atomic model in Figure 1l. Furthermore, the 3D topographic atomic image clearly reveals a lower intensity of Fe compared to Bi atoms, confirming the partial incorporation of Fe into A/C‐Bi₂O₃ (Figure 1m). The asymmetric Fe₁‐O‐Bi electron bridge serves as an activation tunnel that delocalizes the electronic states within Fe₁⊂A/C‐Bi₂O₃, leading to regulated interfacial interactions with LiPSs and significantly enhancing the kinetics of the SRR. Atomic force microscopy (AFM) analysis reveals that the ultra‐thin individual petal‐like structures exhibit thicknesses of 1.66 nm for Fe₁⊂A/C‐Bi₂O₃ and 2.33 nm for C‐Bi₂O₃, respectively. Such atomically thin architectures are conducive to accelerating charge transfer processes, thereby markedly enhancing the efficiency of LiPSs conversion (Figure 1n,o). In addition, energy‐dispersive X‐ray spectroscopy and mapping images further corroborated the uniform distribution of elements across the Fe₁⊂A/C‐Bi₂O₃@CNFs (Figure 1p; Figure S11).

X‐ray absorption spectroscopy was employed to elucidate the chemical state and local coordination environment of Bi atoms (Figure 2a–j). The normalized X‐ray absorption near‐edge structure (XANES) spectra at the Bi L₃‐edge reveal that Fe₁⊂A/C‐Bi₂O₃@CNFs exhibit an evident negative shift in the absorption edge relative to pristine C‐Bi₂O₃@CNFs (Figure 2b). This shift can be attributed to the asymmetric electronic bridges and the robust electronic interaction associated with unsaturated Bi coordination sites at the amorphous–crystalline heterojunction. Linear fitting of the Bi L₃‐edge position further indicates an average Bi oxidation state of approximately +2.3 in Fe₁⊂A/C‐Bi₂O₃@CNFs (Figure 2c). The Fourier‐transform extended X‐ray absorption fine structure (FT‐EXAFS) spectra (Figure 2h) display a prominent coordination shell at ∼2.15 Å, corresponding to Bi─O bonding. Quantitative EXAFS fitting results (Figure 2i,j) demonstrate a pronounced decrease in the Bi coordination number, which is directly associated with the decreased oxidation state of Bi compared to the Bi³⁺ species in C‐Bi₂O₃@CNFs. Wavelet transform (WT) analysis of the Bi L₃‐edge EXAFS signals for both C‐Bi₂O₃@CNFs and Fe₁⊂A/C‐Bi₂O₃@CNFs reveals a dominant scattering contribution centered at ∼6.0 Å⁻¹, characteristic of Bi‐O coordination (Figure 2d–g). Notably, 3D WT contour comparison indicates a substantially higher population of unsaturated Bi coordination number in Fe₁⊂A/C‐Bi₂O₃@CNFs, providing compelling evidence for the formation of Fe₁‐O‐Bi.

FIGURE 2.

(a) Diagram of Synchrotron Radiation. (b) Bi L₃‐edge XANES spectra. (c) Linear relationship between L₃‐edge absorption energy and oxidation state. WT contour plots of C‐Bi₂O₃@CNFs (d, e) and Fe₁⊂A/C‐Bi₂O₃@CNFs (f, g). FT‐EXAFS spectra (h) and the corresponding fitting result for C‐Bi₂O₃@CNFs (i) and Fe₁⊂A/C‐Bi₂O₃@CNFs (j). (k) The total DOS diagrams of Fe₁⊂A/C‐Bi₂O₃@CNFs and C‐Bi₂O₃@CNFs. (l, m, n) The calculated charge density differences of Fe₁⊂A/C‐Bi₂O₃@CNFs. (o) Calculated electrostatic potentials of C‐Bi₂O₃@CNFs and Fe₁⊂A/C‐Bi₂O₃@CNFs. (p) The diagram of electric dipole domains within Fe₁⊂A/C‐Bi₂O₃.

The chemical states of the composites were examined by X‐ray photoelectron spectroscopy (XPS). The high‐resolution Bi 4f spectrum of C‐Bi₂O₃@CNFs could be resolved into two characteristic peaks at 164.6 eV (4f _5/2) and 159.3 eV (4f _7/2), corresponding to Bi³⁺ species (Figure S12). In comparison, the Bi 4f binding energies of Fe₁⊂A/C‐Bi₂O₃@CNFs exhibited an evident shift to lower values, appearing at 164.2 eV (4f _5/2) and 158.9 eV (4f _7/2). This negative shift is mainly ascribed to the formation of asymmetric Fe₁‐O‐Bi electronic bridges and the cooperative effect of the amorphous/crystalline heterostructure, which collectively enhances the local electron density surrounding the Bi atoms. These results provide direct evidence for the electronic structure reconfiguration of Bi species induced by the asymmetric Fe₁‐O‐Bi coordination environment in Fe₁⊂A/C‐Bi₂O₃@CNFs.

The exceptional electrocatalytic performance of Fe₁⊂A/C‐Bi₂O₃ in Li||S batteries can be attributed to its intrinsic amorphous/crystalline heterointerface and the symmetry‐breaking charge distribution. These features work in concert to create spatially directed and decoupled dipole‐active domains, which effectively facilitate SRR. Density of states (DOS) calculations demonstrate that the narrowed bandgap of the Fe₁⊂A/C‐Bi₂O₃, in comparison to C‐Bi₂O₃ (2.3 eV), contributes to an increase in electrical conductivity (Figure 2k). The significantly enhanced electronic conductivity observed in the Fe₁⊂A/C‐Bi₂O₃@CNFs can be primarily attributed to the synergistic interplay between the interfacial built‐in electric field and robust electron delocalization within the amorphous phase. The amorphous phase, characterized by its enhanced electron delocalization, provides a continuous network of pathways featuring lower electronic tunneling barriers and higher intrinsic electron mobility (Figure 2l). This transforms the amorphous domains into efficient electron sponges, capable of rapidly accepting and dispersing the injected electrons, minimizing carrier localization and recombination at defects or interfaces.

The differential charge density analysis (Figure 2m,n) provides definitive evidence of interfacial charge redistribution, characterized by notable electron accumulation and depletion at the crystalline‐amorphous interface. While interfacial polarization is a common feature in heterojunction dipole domains, in the case of Fe₁⊂A/C‐Bi₂O₃@CNFs, it significantly enhances electronic conductivity by facilitating the directional separation and transport of charge carriers. Furthermore, the dipole domain plays a critical role in polarizing the adsorbed polysulfides through dipole‐dipole interactions with LiPSs, effectively lowering the energy barrier for SRR [31, 32].

DFT calculations demonstrate a significant difference in the work function between the crystalline and amorphous phases of Bi₂O₃. As shown in Figure 2o, the work function of C‐Bi₂O₃ (5.1 eV) is higher than that of the amorphous phase (4.3 eV), leading to a spontaneous migration of electrons from the amorphous region to the crystalline region in order to achieve Fermi level alignment. Upon their contact, interfacial charge redistribution occurs, resulting in the formation of an electric dipole domain composed of aligned microscopic dipoles (Figure 2p). This domain establishes a strong built‐in electric field at the macroscopic scale, which facilitates the directional separation and transport of electrons, thereby enhancing overall electronic conductivity. The localized electric field within the dipole domain significantly influences the electrochemical behavior of LiPSs. As polar molecules, LiPSs can experience substantial chemical adsorption, polarization, and cleavage via dipole‐dipole interactions [33]. Electric dipole catalysts often serve as bidirectional catalysts, synergistically facilitating the oxidation and reduction of LiPSs during the charge and discharge processes of Li||S batteries [34]. During the discharge process, the electron‐rich region plays a crucial role in facilitating the reduction of LiPSs, as it provides a surplus of electrons that accelerate the conversion of Li₂S [35]. In parallel, the electron‐deficient region supports the oxidation reactions necessary for charging by efficiently extracting electrons from Li₂S. Moreover, this electric field exerts a directional driving force on Li⁺, effectively reducing the diffusion energy barrier [36].

In Li||S batteries, in situ UV–vis spectroscopy serves as a powerful optical diagnostic tool for dynamically monitoring the dissolution, diffusion, and conversion of LiPSs during electrochemical processes. In this work, in situ UV–vis spectroscopy was utilized to investigate the dynamic evolution of LiPSs mediated by Bi₂O₃‐based catalysts. The measurements were conducted in a customized cell, enabling spectral acquisition at a constant current rate of 0.05 C within the wavelength range of 400−600 nm. The characteristic absorbance at 420 nm, attributed to Li₂S₄, was employed as a key indicator to reflect the dynamic concentration of soluble LiPSs during discharge. Notably, the Fe₁⊂A/C‐Bi₂O₃ electrode exhibited significantly lower UV–vis signal intensity in the Li₂S₄‐related region compared to C‐Bi₂O₃, indicating an effectively suppressed shuttle effect of polysulfide (Figure 3a,b).

FIGURE 3.

Contour maps of in situ UV–vis spectra of the cells with (a) Fe₁⊂A/C‐Bi₂O₃@CNFs and (b) C‐Bi₂O₃@CNFs. (c) Schematic diagram of the in situ XRD setup and corresponding (d) contour maps. (e) Schematic diagram of the in situ Raman setup and corresponding (f) contour maps. (g) UV absorbance spectra of the Li₂S₆ solutions. (h, i) ELF of Fe₁⊂A/C‐Bi₂O₃. (j) The difference charge density of Li₂S₆ adsorbed on Fe₁⊂A/C‐Bi₂O₃ and C‐Bi₂O₃. (k) The PDOS of Bi‐6p orbitals. (l) Schematic illustrations of the electronic interactions among orbitals. (m) Scheme for energy level splitting of Bi‐6p orbitals and S‐3p orbitals. (n) COHP for Bi‐S coupling between LiPSs and Bi‐based catalysts. (o) Adsorption energy between LiPSs and catalysts.

In situ X‐ray diffraction (XRD) was further employed to investigate the phase evolution of sulfur species during discharge, specifically the formation of Li₂S, which is a critical step affecting the reversibility and kinetics of Li||S batteries (Figure 3c). XRD patterns for both C‐Bi₂O₃@CNFs and Fe₁⊂A/C‐Bi₂O₃@CNFs electrodes were collected in real time under a discharge current of 0.1 C. The Fe₁⊂A/C‐Bi₂O₃@CNFs electrode displayed a significantly enhanced diffraction peak corresponding to Li₂S (2θ ≈ 27°, PDF#23‐0369) during the late discharge stage (Figure 3d), while only a weak Li₂S signal was detected in the C‐Bi₂O₃@CNFs‐based cell (Figure S13).

To investigate the evolution of LiPSs during the electrochemical process, in situ Raman analysis was utilized to monitor the LiPSs on the cathode side (Figure 3e). The Raman signals of LiPSs are observed in the range of 150−500 cm⁻¹, resulting from the complex comproportionation and disproportionation processes among various LiPSs species. In C‐Bi₂O₃, Raman signals appear rapidly, indicating a fast and substantial accumulation of unreacted LiPSs (Figure 3f). In stark contrast, the Fe₁⊂A/C‐Bi₂O₃@CNFs electrode demonstrates significantly reduced Raman signal intensities for LiPSs during discharge, indicating effective adsorption and accelerated conversion of LiPSs (Figure S14). The integration of multiple in situ techniques provides compelling evidence that the Fe₁⊂A/C‐Bi₂O₃@CNFs catalyst significantly enhances sulfur redox kinetics and inhibits the LiPSs shuttle effect. The shuttle current test serves as a standardized electrochemical method for quantitatively characterising the intensity of the LiPSs shuttling effect in Li||S batteries [37]. The Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell exhibits a significantly lower shuttling current of approximately 1.1 µA cm⁻², in contrast to C‐Bi₂O₃@CNFs (7.2 µA cm⁻²) and CNFs (14.8 µA cm⁻²), demonstrating its superior ability to inhibit the LiPSs shuttling effect (Figure S15). The pronounced self‐discharge behavior caused by the shuttle effect represents another significant drawback of Li||S batteries. In the case of cells with CNFs and C‐Bi₂O₃@CNFs, the open circuit voltages were significantly lower than those observed in the Fe₁⊂A/C‐Bi₂O₃@CNFs cells. This disparity clearly demonstrates that Fe₁⊂A/C‐Bi₂O₃@CNFs effectively mitigates the self‐discharge behavior in Li||S batteries (Figure S16).

The adsorption capabilities of the LiPSs were visualized by immersing equal amounts of materials in 5 mM Li₂S₆ solutions for 8 h (Figure 3g). Both C‐Bi₂O₃@CNFs and Fe₁⊂A/C‐Bi₂O₃@CNFs exhibited the weakest LiPSs signals in the ultraviolet‐visible absorbance spectra of Li₂S₆, indicating a strong adsorption capacity for LiPSs.

The unique Fe₁‐O‐Bi asymmetric electron bridge in the amorphous/crystalline Bi₂O₃ heterojunction acts as a highly effective active site for mediating LiPSs adsorption and conversion in Li||S batteries. The inherent differences in atomic radius, electronic configuration, and electronegativity between Fe and Bi atoms introduce an asymmetric electronic configuration along the Fe₁‐O‐Bi bridge. This elemental disparity disrupts the potential inversion symmetry that could exist in the symmetrical Bi‐O‐Bi bridge, resulting in a polarized electronic state. Electronic localization function (ELF) analysis confirms substantial electron accumulation around the Bi sites of the Fe₁‐O‐Bi structure within Fe₁⊂A/C‐Bi₂O₃, which arises directly from the electronic redistribution induced by the breaking of symmetry along the asymmetric Fe₁‐O‐Bi electron bridge (Figure 3h,i). Specifically, the strong electron‐donating nature of Fe leads to excessive electron density transfer to the bridging oxygen, which subsequently weak the electron donation from the adjacent Bi atom, resulting in a localized electron‐rich state at the Bi site (Figure 3l). This electronic redistribution directly influences the electronic structure of the Bi active center, causing a downshift of Bi 6p‐band center (Figure 3k). The lowered p‐band center optimizes the interaction between Bi and S, resulting in a moderate adsorption strength for LiPSs by increasing electron occupation in the Bi─S anti‐bonding orbitals, which prevents overly strong binding that could passivate the active sites (Figure 3m). Figure 3j and Figures S17 and S18 exhibit the analysis of the charge density difference of adsorbed LiPSs on Fe₁⊂A/C‐Bi₂O₃ and C‐Bi₂O₃. The results indicate that the electron transfer from Fe₁⊂A/C‐Bi₂O₃ to LiPSs is lower than that from C‐Bi₂O₃, indicating that Fe₁⊂A/C‐Bi₂O₃ could facilitate the dynamic chemisorption‐desorption of LiPSs and enhance the reversible redox reaction of LiPSs. The interfacial interactions between LiPSs and Fe₁⊂A/C‐Bi₂O₃@CNFs were further investigated by XPS (Figure S19). Compared with pristine Li₂S₆, the Li 1s peak of Fe₁⊂A/C‐Bi₂O₃@CNFs displayed a pronounced shift toward higher binding energy. The Fe 2p spectra of Fe₁⊂A/C‐Bi₂O₃@CNFs after interaction with Li₂S₆ displayed an evident shift toward lower binding energies. This downshift signifies electron donation from LiPSs to the Fe active sites, confirming that the Fe atoms within the Fe₁‐O‐Bi electronic bridges act as key catalytic sites that effectively suppress the shuttle effect of LiPSs.

According to the Sabatier principle, the interaction between catalyst and reactant molecules must be finely balanced, neither excessively strong nor too weak, to attain optimal catalytic performance. In Li||S batteries, the sulfur host composites should exhibit moderate adsorption capability toward LiPSs. This moderated interaction facilitates the reversible adsorption and desorption of LiPSs, accelerating the redox kinetics of sulfur. To further clarify the influence of Bi‐S coupling on the adsorption system, projected Crystal Orbital Hamilton Population (COHP) analyses were conducted. As illustrated in Figure 3n, for intermediates ranging from Li₂S₈ to Li₂S₄, a portion of the anti‐bonding states associated with the Bi─S bond in both C‐Bi₂O₃ and Fe₁⊂A/C‐Bi₂O₃ shifts above the Fermi level. The integrated −COHP (−ICOHP) values of Bi‐S for Fe₁⊂A/C‐Bi₂O₃ are generally lower than those of C‐Bi₂O₃, indicating an moderate electronic coupling between Bi and S. These findings are further corroborated by the calculated adsorption energies, as shown in Figure 3o.

Li||S batteries with a sulfur loading of 1.5 mg cm⁻² were assembled to investigate the differences in electrochemical performance based on different catalysts. Figure S20 delivers cyclic voltammetry (CV) curves within the voltage window of 1.7−2.8 V under the sweeping rate of 0.1 mV s⁻¹. Two prominent negative peaks occurred at about 2.31 V (peak B) and 2.03 V (peak C), which are associated with the reduction of S₈ to high‐order soluble Li₂S_n (4 ≤ n ≤ 8) and followed by the reduction to low‐order insoluble Li₂S₂/Li₂S, respectively. Additionally, two close positive peaks located at 2.30 and 2.39 V (Peak A) are related to the oxidation of short‐chain Li₂S₂/Li₂S to long‐chain Li₂S_n and finally to S₈, respectively. Furthermore, the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell exhibited the highest peak currents compared to those of C‐Bi₂O₃@CNFs and CNFs, indicating its effective electrocatalytic activity for the conversion of LiPSs (Figure S21). More importantly, as shown in Figure S22, the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell demonstrated the highest onset potential for negative peaks (peaks C) and the lowest onset potential for positive peaks (peak A). This finding confirms the accelerated reaction kinetics of the sulfur redox reaction facilitated by Fe₁⊂A/C‐Bi₂O₃@CNFs, significantly reducing the polarization during the liquid‐solid conversion.

To further investigate the electrochemical performance, CV tests were conducted at different scan rates to evaluate Li⁺ diffusion in Li||S batteries. As illustrated in Figure 4a–d and Figures S23 and S24, the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell exhibits significantly enhanced current densities across a range of scanning rates compared to C‐Bi₂O₃@CNFs and CNFs. The diffusion coefficient of Li⁺ can be qualitatively evaluated using the Randles‐Sevcik equation [38]. The Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell demonstrates the highest Li⁺ diffusivity, as evidenced by the steeper slope observed in the linear fitting of Figure 4b,d, in comparison to the C‐Bi₂O₃@CNFs and pristine CNFs. Additionally, DFT calculations confirm that Fe₁⊂A/C‐Bi₂O₃@CNFs could significantly facilitate Li⁺ diffusion more than C‐Bi₂O₃. Compared with C‐Bi₂O₃ (0.87 eV), the Li⁺ diffusion barrier along the diffusion coordinate on the Fe₁⊂A/C‐Bi₂O₃ (0.46 eV) is lower (Figure 4e,f). The enhancement of Li⁺ diffusion can be attributed to the electronic dipole region, which facilitates flexible, low‐energy migration pathways through a disordered network of under‐coordinated atoms.

FIGURE 4.

Contour plots of (a) peak A and (c) peak C from the CV curves of Li||S batteries. The linear relationships between I _p and the square root of scan rates for (b) peak A and (d) peak C. The energy barrier for the diffusion process of Li⁺ within (e) C‐Bi₂O₃ and (f) Fe₁⊂A/C‐Bi₂O₃. (g) LSV curves and corresponding (h) Tafel plots. Contour plots of CV patterns for (i) Fe₁⊂A/C‐Bi₂O₃@CNFs, (j) C‐Bi₂O₃@CNFs, and (k) CNFs at different temperatures. (l) The linear relationship between ln (I _p) and 1/T. (m) Tafel plots, (n) EIS plots, and (o) chronoamperometric curves for the symmetric cells. (p) DRT curves of the Li||S battery. EIS plots and corresponding DRT contour plots at different temperatures for (q) Fe₁⊂A/C‐Bi₂O₃@CNFs, (r) C‐Bi₂O₃@CNFs, and (s) CNFs‐based cells. (t) The linear relationship between the ln(1/R) and the 1/T.

The oxidation behavior of Li₂S on Fe₁⊂A/C‐Bi₂O₃ was systematically investigated using three‐electrode linear sweep voltammetry (LSV). In this setup, a 0.1 m Li₂S/methanol solution served as the electrolyte, while platinum and Ag/AgCl electrodes functioned as the counter and reference electrodes, respectively. Notably, Fe₁⊂A/C‐Bi₂O₃ exhibited the lowest onset voltage and the highest current response compared to C‐Bi₂O₃ (Figure 4g). In addition, the Tafel plots of the LSV revealed that Fe₁⊂A/C‐Bi₂O₃ displays the smallest slope of 95 mV dec⁻¹, compared to 102 mV dec⁻¹ for C‐Bi₂O₃ and 119 mV dec⁻¹ for CNFs (Figure 4h). These results underscore the enhanced electrocatalytic performance of Fe₁⊂A/C‐Bi₂O₃ in Li||S batteries.

The activation energy (E _a) is used to deepen the understanding of the role of Fe₁⊂A/C‐Bi₂O₃ in facilitating the conversion of LiPSs. Temperature‐dependent CV measurements were performed at varying temperatures (Figure 4i–k and Figure S25). Given that the conversion of LiPSs to solid Li₂S is the rate‐determining step, particular attention was paid to calculating E _a for this liquid‐to‐solid transformation. According to the Arrhenius equation, the peak current is proportional to the reaction rate. Linear fitting of the current versus 1/T revealed that the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell exhibited the smallest slope, indicating that Fe₁⊂A/C‐Bi₂O₃@CNFs effectively lowers the activation energy barrier for the nucleation and deposition of Li₂S (Figure 4l). As illustrated in Figure 4m, the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell exhibited the highest response current during both anodic and cathodic processes, alongside the smallest Tafel slope and the highest exchange current density of 2.5 × 10⁻⁴ mA cm⁻², significantly surpassing that of C‐Bi₂O₃@CNFs (3.2 × 10⁻⁶ mA cm⁻²). These results demonstrate that Fe₁⊂A/C‐Bi₂O₃@CNFs substantially enhances the redox kinetics of LiPSs, effectively improving the electrocatalytic performance of Li||S batteries.

CV curves of Li₂S₆ symmetric batteries were further utilized to evaluate the electrocatalytic activity of the Fe₁⊂A/C‐Bi₂O₃@CNFs. As shown in Figure S26, the Fe₁⊂A/C‐Bi₂O₃@CNFs electrode demonstrates a larger peak area and a higher peak current response compared to both the C‐Bi₂O₃@CNFs and CNFs electrodes, confirming the enhanced redox kinetics. Moreover, the redox peaks of the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell maintain a stable shape with only slight shifts and narrow peak separations, even at elevated scan rates, indicating exceptional electrochemical reversibility. In contrast, the CV profiles of the C‐Bi₂O₃@CNFs and CNFs electrodes exhibit a narrow and elongated form, lacking discernible redox peaks (Figure S26b). Notably, the electrochemical impedance spectroscopy (EIS) plots for the Fe₁⊂A/C‐Bi₂O₃@CNFs symmetric cells reveal a smaller charge transfer resistance (R _ct) compared to those of the C‐Bi₂O₃@CNFs and CNFs electrodes (Figure 4n). In addition, chronoamperometry curves for Fe₁⊂A/C‐Bi₂O₃@CNFs showed an enhanced current response compared to C‐Bi₂O₃@CNFs and CNFs, highlighting the beneficial effect of accelerating the redox reaction of LiPSs (Figure 4o; Figure S27).

Figure 4p illustrates that the impedance of the Li||S batteries can be effectively analyzed by segmenting it into distinct regions through distribution of relaxation times (DRT), yielding valuable insights into inter‐particle ohmic resistance (P₁), charge transfer at the positive electrode (P_2‐4), and diffusion of LiPSs (P₅). The performance of Li||S batteries utilizing Fe₁⊂A/C‐Bi₂O₃@CNFs and C‐Bi₂O₃@CNFs was evaluated across a range of temperatures, complemented by corresponding DRT analysis. The Nyquist plots presented in Figure 4q–s reveal the presence of semicircles and linear tails, where the diameter of the semicircle is typically indicative of the R _ct associated with the conversion of LiPSs. As the temperature increases, the Fe₁⊂A/C‐Bi₂O₃@CNFs electrode exhibits a continuous reduction in ohmic resistance, as well as in both the LiPSs‐related diffusion resistance and the charge‐transfer resistance linked to the SRR. This trend collectively signifies a marked enhancement in the kinetics of the SRR. Furthermore, throughout the entire temperature range investigated, the DRT spectra of Fe₁⊂A/C‐Bi₂O₃@CNFs consistently display lower intensity peaks corresponding to LiPSs diffusion and electrochemical conversion processes. In addition, the charge‐transfer resistance at different temperatures shows a linear relationship between the inverse of the absolute temperature and the logarithm of the reciprocal of the charge‐transfer resistance (Figure 4t). The derived slope from the linear fit indicates that the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell exhibits the lowest activation energy (E _a), suggesting that Fe₁⊂A/C‐Bi₂O₃@CNFs effectively reduces the energy barrier of SRR.

As illustrated in Figure 5a,b, in situ EIS was utilized to monitor the charge transfer resistance during the discharging process. Notably, when the discharge voltage declines to 1.7 V, there is a significant increase in transfer resistance, resulting from the formation of insulating solid‐phase Li₂S₂/Li₂S and the sluggish kinetics of the rate‐determining step. To quantitatively analyze the sulfur reduction kinetics during discharge, in situ DRT analysis was performed. This method enables the deconvolution of polarization contributions from each individual electrode process. Notably, the DRT results indicate that Fe₁⊂A/C‐Bi₂O₃@CNFs significantly reduce both the diffusion resistance of LiPSs (−0.5< log τ < 0.5) and the charge transfer resistance at the positive electrode (−2< log τ < −1) compared to C‐Bi₂O₃@CNFs. This result is closely associated with the accelerated reaction rate of the SRR during the discharge process, which decreases the viscosity of the electrolyte and facilitates the diffusion of both Li⁺ and LiPSs.

FIGURE 5.

(a, b) In situ EIS and the corresponding DRT contour of Li||S cells during the discharging process. (c) Kinetic assessment using PITT measurement. (d, f) Dimensionless current‐time transient profile during the deposition of Li₂S. (e) Rate performance of the Li||S batteries. (g) Galvanostatic discharge profiles of Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell. (h) Diagram of lithium metal corrosion caused by the shuttle effect of LiPSs. (i, j) AFM images of the lithium anode after cycling. (k) GITT curves of Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cells. (l, m, n) COHP of Li‐S coupling in Li₂S. (o, p) Energy barrier of Li₂S decomposition.

To investigate the kinetics of the liquid‐solid conversion during the discharging process, the potentiostatic intermittent titration technique (PITT) was employed on cells with a lithium anode, an electrocatalyst‐coated cathode, and a Li₂S₈ catholyte (Figure 5c; Figure S28a,b). At higher titration potentials above 2.09 V, long‐chain LiPSs were first electrochemically reduced to short‐chain LiPSs. Subsequently, at a lower titration potential of 2.09 V, the soluble short‐chain LiPSs were further reduced to solid Li₂S, resulting in deposition on the surface of the electrocatalysts. The Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell reached the peak current in the shortest time and achieved the highest maximum current response (0.78 mA), surpassing those of C‐Bi₂O₃@CNFs (0.64 mA) and CNFs (0.37 mA). Additionally, the deposition of Li₂S on Fe₁⊂A/C‐Bi₂O₃@CNFs (526.91 mAh g⁻¹) was significantly higher than that on C‐Bi₂O₃@CNFs (371.09 mAh g⁻¹) and CNFs (260.24 mAh g⁻¹), respectively. The dimensionless current‐time transients were further plotted to elucidate the Li₂S deposition mode on different electrocatalysts (Figure 5d,f; Figure S28c). The C‐Bi₂O₃@CNFs electrode exhibits the conventional 2D Li₂S deposition mode. In this process, the catalytic active sites were obscured beneath the insulating Li₂S, inhibiting the continuous deposition of Li₂S. In contrast, the Li₂S deposition on Fe₁⊂A/C‐Bi₂O₃@CNFs closely followed a three‐dimensional instantaneous (3DI) deposition model, resulting in a higher deposition capacity (Figure 5d).

As depicted in the galvanostatic charge‐discharge profiles (Figure S29a), the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell achieves an impressive initial discharge capacity of 1422 mAh g⁻¹ at a current density of 0.2 C, which markedly exceeds those of the C‐Bi₂O₃@CNFs (1198 mAh g⁻¹) and CNFs (933 mAh g⁻¹). As presented in Figure S29a,b, the ratio of capacities associated with the two discharge platforms (denoted as Q_L/Q_H) reflects the varying electrocatalytic activity of LiPSs during the discharge process. The phase conversion coefficient of Fe₁⊂A/C‐Bi₂O₃@CNFs (2.26) is notably higher than that of C‐Bi₂O₃@CNFs (2.14) and CNFs (2.10), indicating a superior sulfur utilization capability and more effective conversion of LiPSs throughout the discharge process. Additionally, as shown in Figure S29c, the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell demonstrates the lowest polarization (ΔE = 188 mV) compared to the C‐Bi₂O₃@CNFs (205 mV) and the CNFs (225 mV). Furthermore, a distinct valley is observed between the high plateau and the low plateau, referred to as the “nucleation point” (Figure S29d–f). The potential difference between the nucleation point and the tangential line of the potential plateau serves as an indicator for evaluating the nucleation kinetics of Li₂S. The CNFs‐based cell displays the largest overpotential of 24 mV, indicating the presence of a significant interfacial energy barrier for Li₂S nucleation and deposition. In contrast, the nucleation behavior of Li₂S in the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell is markedly different, exhibiting a substantially reduced interfacial energy barrier of only 11 mV. The rate capabilities of the Li||S batteries were further evaluated. As depicted in Figure 5e, the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell demonstrated excellent rate performance, achieving high capacities of 1428, 1083, 977, and 660 mAh g⁻¹ at current densities of 0.2, 0.5, 1.0, and 5.0 C, respectively. Notably, when the current was returned to 0.2 C, the reversible capacity was restored to 1283 mAh g⁻¹, which corresponds to 90% of the initial capacity.

As shown in Figure 5g and Figure S30, the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell maintained a capacity of 1165 mAh g⁻¹ after 200 cycles, with a decay rate of just 0.08% per cycle. This performance significantly outperforms that of C‐Bi₂O₃ (842 mAh g⁻¹, 0.15% per cycle) and CNFs (347 mAh g⁻¹, 0.32% per cycle). Furthermore, at a high current density of 1.0 C, the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell delivered an initial capacity of 945 mAh g⁻¹ and maintained a capacity of 698 mAh g⁻¹ after 1000 cycles, resulting in a capacity retention of 74% and an average degradation rate of 0.026% per cycle (Figure S31). The shuttle effect of LiPSs leads to substantial corrosion of the lithium metal anode, adversely affecting both electronic conductivity and the kinetics of electrochemical reactions (Figure 5h). The AFM images and digital photographs of the lithium metal anode assembled in the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell after cycling reveal a relatively smooth surface compared to that of the C‐Bi₂O₃@CNFs. This suggests that Fe₁⊂A/C‐Bi₂O₃@CNFs can effectively adsorb LiPSs, promote uniform lithium deposition, and mitigate lithium metal corrosion. (Figure 5i,j; Figure S32). Galvanostatic intermittent titration technique (GITT) was further employed to evaluate the influence of Fe₁⊂A/C‐Bi₂O₃@CNFs on the internal resistance during the discharge process of Li||S batteries (Figure 5k; Figure S33). Notably, the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell demonstrates the lowest internal resistance for both the nucleation and deposition of Li₂S when compared to the C‐Bi₂O₃@CNFs and CNFs, highlighting the enhanced kinetics of the sulfur reduction reaction.

The COHP analysis was employed to quantitatively assess the Li‐S coupling interactions of Li₂S. The integrated COHP values down to the Fermi level (ICOHP) for the Li─S bonds in these adsorption systems were calculated to evaluate bonding strength. A lower ‐ICOHP value indicates a weaker interaction between the lithium and sulfur atoms. As shown in Figure 5l–n, the ICOHP for the Li─S bond in the Fe₁⊂A/C‐Bi₂O₃ system is notably lower than that in the C‐Bi₂O₃ system, indicating that Fe₁⊂A/C‐Bi₂O₃ can effectively polarize and facilitate the cleavage of the Li─S bond. This phenomenon can be attributed to dipole‐dipole interactions mediated by the electric dipole domain. The weakened Li─S bond notably decreases the energy barrier for the oxidation of Li₂S during the charging process, aligning with the calculations related to the decomposition energy barrier of Li₂S. As demonstrated in Figure 5o,p, the energy barrier for Li₂S decomposition on Fe₁⊂A/C‐Bi₂O₃ is 0.76 eV, significantly lower than the 1.36 eV observed for C‐Bi₂O₃, confirming the more favorable kinetics for Li₂S oxidation during the initial stages of the charging process.

Considering practical applications, high‐sulfur‐loading Li||S batteries were fabricated. As depicted in Figure 6a, two distinct voltage plateaus remain clearly visible even under a high sulfur loading of 7.4 mg cm⁻². The desirable cycle stability was acquired under 0.2 C, with high capacity retentions of 4.9 and 4.6 mAh cm⁻² after 200 cycles under sulfur loadings of 7.4 and 5.2 mg cm⁻², respectively (Figure 6b). Obviously, the acquired areal capacity under high sulfur loading is competitive with recent publications (Figure 6c and Table S1) [39, 40].

FIGURE 6.

(a) Discharge profiles and (b) cycling performance of Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cells with high sulfur loadings. (c) Performance comparisons of Li||S batteries with recent references. (d) Fabrication process of the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based pouch cell. (e) Cycling performance of Li||S pouch cell. (f) Performance comparison of the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based pouch cell with recent literature. (g) Practical utilization of the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based pouch cell. (h) Schematic illustration of the enhanced SRR kinetics mediated by the Fe₁⊂A/C‐Bi₂O₃. (i–k) Gibbs free energy for SRR.

A significant amount of black liquor lignin is generated during the pulp and paper production process. In this work, lignin was employed to fabricate high‐performance CNFs via electrospinning for constructing high‐performance Li||S pouch cells (Figure 6d). By carefully controlling the weight ratios of each key component, a Li||S pouch cell with a total weight of 5.8 g and a high energy density of 387 Wh kg⁻¹ was successfully fabricated (Table S2). The 1.2 Ah Fe₁⊂A/C‐Bi₂O₃@CNFs‐based Li||S pouch cell maintains a capacity retention of 83% after 50 cycles at 0.05 C (Figure 6e). Compared to the reported Ah‐class Li||S pouch cells [41, 42, 43, 44, 45, 46, 47, 48, 49], the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell exhibits significant advancements in total capacity (1.2 Ah) and weight energy density (387 Wh kg⁻¹) at a relatively high sulfur loading of 10 mg cm⁻² and an extremely low E/S ratio of 2.8 µL mg⁻¹ (Figure 6f). Additionally, Figure 6g demonstrates the practical application of the Fe₁⊂A/C‐Bi₂O₃@CNFs‐based cell, successfully serving as a power source to charge mobile phones and drive drone flights, highlighting its robust performance and potential for real‐world energy storage applications.

The preferable Li||S electrochemical performance mediated with the Fe₁⊂A/C‐Bi₂O₃@CNFs was illustrated in Figure 6h. The superior electrocatalytic performance of Fe₁⊂A/C‐Bi₂O₃@CNFs arises from the synergistic catalytic effects of its amorphous structure, electronic dipole domains, and the asymmetric Fe₁‐O‐Bi electronic bridge with symmetry breaking. This synergy leads to an optimized electronic structure that lowers the p‐band center and moderates the adsorption of LiPSs, facilitating dynamic chemical adsorption and desorption processes. Consequently, Fe₁⊂A/C‐Bi₂O₃@CNFs can effectively reduce reaction energy barriers and contribute to the high performance of Li||S batteries. This was further substantiated by the decrease in Gibbs free energy difference (ΔG), as illustrated in Figure 6i–k. Among the various multistep SRRs, the nucleation of Li₂S/Li₂S₂ exhibited the highest energy barrier, thereby serving as the rate‐determining step. Significantly, the Fe₁⊂A/C‐Bi₂O₃ exhibited a markedly lower energy barrier compared to C‐Bi₂O₃, indicating that Fe₁⊂A/C‐Bi₂O₃ can effectively reduce the energy barrier for the reduction of LiPSs.

3. Conclusion

In summary, this study presents an ingenious design paradigm for a high‐performance Li||S battery by leveraging the synergistic effects of amorphization‐triggered symmetry breaking and electric dipole domain. The amorphous phase, in conjunction with the asymmetric Fe₁‐O‐Bi electronic bridges, facilitates electron delocalization, reducing the p‐band center of the Bi active sites. The electronic dipole domain at the heterointerface of Fe₁⊂A/C‐Bi₂O₃ creates rapid transport channels for charge carriers and Li⁺, effectively mitigating polarization of Li||S batteries. Furthermore, the strong dipole‐dipole interactions at the electronic dipole domain interface destabilize the LiPSs molecular configuration, reducing the energy barrier for the cleavage of Li─S/S─S bonds. As a consequence, Fe₁⊂A/C‐Bi₂O₃@CNFs exhibits exceptional electrochemical performance and significant practical application. An Ah‐level pouch cell (1.2 Ah) was successfully developed, achieving an energy density of 387 Wh kg⁻¹, even under high sulfur loadings of 10 mg cm⁻² with a lean electrolyte (E/S = 2.8 µL mg⁻¹). This work elucidates the critical link between the symmetry‐breaking electronic configuration triggered by amorphization and the enhanced electrochemical performance, offering a new design strategy for practical Li||S batteries.

Conflicts of Interest

The authors declare no conflict of interest.

Supporting information

Supporting File: adma73469‐sup‐0001‐SuppMat.docx,

Data Availability Statement

The data that support the findings of this study are availablen the supplementary material of this article.