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. 2026 Jun 30. Online ahead of print. doi: 10.1039/d6sc04034a

Breaking the strength–dendrite paradox in polymer electrolytes: spherical lithium deposition via redox-active Fe–O/Cl centers

Ting Zhang a,†, Yifan Xu a,b,†, Zijian He a, Junliang Zhao c, Lei Gao d, Songbai Han b, Chenghua Sun e, Qiang Xu f, Ruo Zhao a,✉
PMCID: PMC13344480  PMID: 42422098

Abstract

Polyethylene oxide (PEO)-based electrolytes are essential for solid-state batteries due to their intrinsic flexibility and affordability for large-scale manufacturing. However, they suffer from severe Li dendrite growth. Although traditional strategies focus on improving the mechanical modulus to resist dendrites, they lead to compromised interfacial compatibility. Herein, we propose a cation–anion synergy strategy to regulate Li plating behavior and achieve spherical Li deposition in PEO-based electrolytes. Specifically, Fe–O/Cl centers are formed by simply incorporating ferric chloride (FeCl3) into a PEO matrix, which serve as nucleation sites for Li plating. The electronegative Cl− anions act as a “Li+ pump” to concentrate Li+ around nucleation centers, whereas redox-active Fen+ cations function as an “electron reservoir” that traps interfacial electrons and enables preferential Li deposition. Meanwhile, Fen+ cations catalyze TFSI− decomposition, constructing a robust LiF/Li2O-rich interfacial layer. This collaborative cation–anion synergy enables highly uniform spherical Li plating. Consequently, the as-prepared FeCl3-modified electrolyte with a lower modulus (3.1 MPa) exhibits superior ability against dendrite growth, breaking the conventional strength–dendrite paradox. The obtained electrolyte enables a high critical current density of 2.8 mA cm−2. The symmetric cell sustains stable cycling over 3000 h, while the full cell maintains a high capacity of 129.4 mAh g−1 after 1000 cycles. This work provides a facile and effective route to regulate Li deposition for high-performance polymer electrolytes.

FeCl3-doped polyethylene oxide (PEO) forms uniformly dispersed Fe–O/Cl centers. Electronegative Cl− acts as a “Li± pump” and Fen+ as an “electron reservoir” to enrich Li+ flux, inducing spherical Li deposition for long-cycle batteries.graphic file with name d6sc04034a-ga.jpg

1. Introduction

Li-metal-based solid-state batteries (LMSSBs) are poised to revolutionize energy storage technologies by offering significantly higher energy density compared to conventional liquid Li-ion batteries.1–3 Solid-state electrolytes (SSEs) are the key component and appear as a promising way to pair with Li anodes since they are solvent-free and have no leakage risk, enhancing battery safety.4–6 Among various SSEs, composite polymer electrolytes, particularly those based on polyethylene oxide (PEO), stand out for their intrinsic flexibility, low cost, and excellent compatibility with industrial roll-to-roll manufacturing.7,8 In recent years, there have been significant improvements in the ionic conductivity, mechanical strength, and electrochemical stability of PEO-based electrolytes by adding solid fillers,9–15 plasticizers,16–19etc. However, the fundamental issue of Li dendrite growth has not yet been resolved. These dendrites can pierce polymer electrolytes, triggering internal short circuits and even severe thermal runaway, which is the primary barrier to the practical application of PEO-based electrolytes.

Previous studies predominantly ascribe dendrite penetration to the low elastic modulus of the PEO matrix, which cannot withstand the interfacial mechanical stress generated by inhomogeneous Li plating. This viewpoint has driven extensive research into high-strength composite electrolytes, such as fiber-reinforced composites and multilayer architectures.20–22 However, these strategies inevitably face trade-offs, including compromised flexibility, degraded interfacial compatibility and complicated manufacturing processes. This gives rise to the strength–dendrite paradox, where mechanical suppression alone cannot address the root cause of Li dendrite propagation. Rather than solely pursuing higher modulus, recent insights suggest that regulating Li nucleation behavior and deposition morphology holds the key to resolving this long-standing issue. In particular, spherical Li deposits can greatly mitigate interfacial stress concentration, eliminating the driving force for dendrite penetration. Thus, there is an urgent need for strategies that move beyond mechanical reinforcement to achieve precise regulation of Li deposition behavior for dendrite-free PEO-based electrolytes.

The morphology of deposited Li is determined by its nucleation and growth kinetics, which involve coupling of Li+ migration, interfacial charge transfer, and electro-crystallization.20,23 The classic space charge theory proposed by Chazalviel et al. demonstrates that local Li+ depletion at the interface is the primary trigger for dendrite growth.24,25 Meanwhile, the well-known tip effect shows that randomly formed small-size Li nuclei at interfaces act as charge accumulation sites, driving preferential Li+ migration and longitudinal growth into high-aspect-ratio whisker-like or moss-like dendrites.21 These dendrites protrude throughout electrolytes and cause battery failure.22,23 Furthermore, classical nucleation theory and Butler–Volmer relations reveal that the initial nucleus size is inversely proportional to the nucleation overpotential and exchange current density.26–29 Lower overpotential enables the formation of larger and more uniform Li nuclei. Therefore, constructing Li+-philic interfaces with uniform nucleation sites is an effective strategy to compensate for local Li+ exhaustion, reduce the nucleation barrier, promote uniform Li nucleus formation, and ultimately achieve dendrite-free Li deposition.

In this work, we report a cation–anion synergy strategy to realize highly uniform spherical Li nucleation in PEO-based electrolytes, thus achieving a dendrite-free interface and long-lifespan batteries. Mechanistic studies further reveal the interplay of redox-active Fe–O/Cl centers in governing Li deposition behavior (Scheme 1). Specifically, Fe–O/Cl centers were formed by simply incorporating ferric chloride (FeCl3) into the PEO matrix, which serve as nucleation sites for Li plating. The electronegative Cl− anions with high Li+ affinity act as a “Li+ pump” to enrich Li+ flux and accelerate directional Li+ migration toward the nucleation centers. Meanwhile, the redox-active Fen+ cations with high electron-withdrawing ability serve as an “electron reservoir” to localize interfacial electrons, enabling preferential Li+ reduction at the designed nucleation sites. This collaborative cation–anion synergy enables highly uniform spherical Li plating. Moreover, the Fen+ species catalyzed TFSI− decomposition, forming a robust LiF/Li2O-rich interphase, which favored lateral Li plating.30 Benefiting from these merits, the as-prepared FeCl3-modified electrolyte with a lower modulus exhibits superior ability against dendrite growth, breaking the conventional strength–dendrite paradox. The symmetric cells can run stably over 3000 hours, while the full cell delivers exceptional cycling stability with a capacity of 129.4 mAh g−1 after 1000 cycles. This work provides new insights and facile strategies for regulating Li deposition behavior.

Scheme 1. Interplay of redox-active Fe–O/Cl centers for spherical Li deposition.

Scheme 1

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2. Results and discussion

2.1. Synthesis and investigation of FeCl3-incorporated PEO-based electrolytes

A simplified baseline system consisting of PEO and LiTFSI (denoted as PEO-LiTFSI) was employed as the reference matrix. Then, a small amount (3 wt%) of ferric chloride (FeCl3) was incorporated into the base matrix, which was denoted as FeCl3-PEO-LiTFSI (Fig. S1). X-ray diffraction (XRD) patterns revealed characteristic peaks that were ascribed to crystalline PEO (Fig. 1a). An upward shift towards high diffraction angles was observed after introducing FeCl3, suggesting a decreased interplanar distance. There were no XRD peaks assigned to FeCl3, indicating its complete dissociation and formation of a coordination complex. The Raman spectrum showed a new peak at 330 cm−1 for FeCl3-PEO-LiTFSI, which could be attributed to Fe–O bonds (Fig. 1b).31 X-ray photoelectron spectroscopy (XPS) was further conducted to investigate the chemical interactions (Fig. 1c and d). In the O 1s spectrum of PEO-LiTFSI, the peak at 533.0 eV can be attributed to the –C–O– bonds. In contrast, the O 1s spectrum of FeCl3-PEO-LiTFSI revealed a new peak at 530.8 eV, which was assigned to Fe–O bonds.31,32 The Fe 2p XPS spectrum of FeCl3-PEO-LiTFSI showed four main peaks at 709.3 and 722.6 eV, attributed to Fe2+, and the peaks at 711.8 and 724.2 eV, assigned to Fe3+.33,34 These peaks confirm the coexistence of Fe2+ and Fe3+ oxidation states, corresponding to Fe–O and Fe–Cl bonds. This variable valence state of Fen+ ions reflects their redox properties, establishing a base for electron uptake/release during Li plating.

Fig. 1. (a) XRD patterns, (b) Raman spectra, and (c) O 1s spectra of FeCl3-PEO-LiTFSI and PEO-LiTFSI electrolytes. (d) Fe 2p spectrum of FeCl3-PEO-LiTFSI. (e) Fe K-edge XANES spectra, inset shows the enlarged view. (f) Magnitude of k3-weighted Fourier transforms of the EXAFS spectra. (g) Fitting analysis of FeCl3-PEO-LiTFSI in R space. (h) Wavelet transformation analysis from EXAFS spectra. (i) Calculated structure from DFT analysis.

Fig. 1

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The detailed coordination structure was studied using X-ray absorption near-edge structure (XANES) and extended X-ray absorption fine structure (EXAFS) spectra. The Fe K-edge XANES spectrum of FeCl3-PEO-LiTFSI displayed the absorption edge position between Fe and Fe2O3, revealing that the valence state of Fe species was between Fe and Fe3+ (Fig. 1e). The Fourier-transformed Fe K-edge EXAFS spectra of FeCl3-PEO-LiTFSI exhibited a dominant peak at around 1.75 Å, which can be assigned to Fe–O/Fe–Cl coordination (Fig. 1f).32,35 The EXAFS curve-fitting analysis revealed coordination numbers of 0.6 ± 0.2 for Fe–O and 2.1 ± 0.1 for Fe–Cl, respectively, confirming the presence of Fe–O/Cl coordination centers (Fig. 1g and h). DFT calculations further confirmed the Fe–O and Fe–Cl interactions, revealing that the PEO chain can coordinate with Fe3+ and form Fe–O/Cl centers (Fig. 1i). Notably, the unique Fe–O/Cl centers contain both redox-active Fen+ and electronegative Cl−, which work synergistically to regulate Li plating.

The effect of FeCl3 contents on electrolyte properties were first investigated and optimized. Specifically, the thermal stability of the FeCl3-doped electrolytes exhibited a decreasing trend with increasing FeCl3 content (Fig. S2), which is attributed to the catalytic role of Fe3+ in accelerating polymer degradation. Moreover, the slurry viscosity and film-forming ability of doped systems deteriorated significantly as the FeCl3 loading increased. At a doping level of 4 wt%, the formation of intact and self-supporting films was no longer feasible, thus rendering this concentration unsuitable for further investigation. Differential scanning calorimetry (DSC) analysis revealed that, compared with PEO-LiTFSI, the FeCl3-doped electrolytes exhibit a slight decrease in the melting temperature (Tm) from 51.5 to 48.7 °C and an increase in the glass transition temperature (Tg) from −42.5 to −41.7 °C (Fig. 2a and S3). The decrease in Tm suggests reduced PEO crystallinity and a higher proportion of amorphous domains, which facilitate ion migration. Meanwhile, the slight Tg elevation may arise from local coordination between Fe3+ and ether oxygen in the PEO backbone. Importantly, this interaction also weakens Li+–PEO binding affinity, promoting faster Li+ hopping between adjacent coordination sites. Consequently, the ionic conductivity increased progressively to 3.40, 7.11, and 16.60 × 10−4 S cm−1 at 60 °C with 1, 2, and 3 wt% FeCl3, respectively (Fig. 2b, c, S4 and S5). In parallel, the Li+ transference number also improved slightly with increasing Fe3+ content (Fig. 2d, e and S6), which was ascribed to the interaction between Lewis acidic Fe3+ and TFSI− anions. Moreover, Fe–O/Cl centers stabilized polymer chains and mitigated its oxidative decomposition, thus improving the electrochemical stability window (Fig. 2f).

Fig. 2. (a) DSC curves, (b) Nyquist plots, and (c) Arrhenius plots of PEO-LiTFSI and FeCl3-PEO-LiTFSI. (d) Li+ transference number, inset shows the Nyquist plots before and after polarization. (e) Comparative bar chart. (f) Linear sweep voltammetry (LSV) curves (g) AFM DMT modulus maps, and (h) nanoindentation load–displacement curves of PEO-LiTFSI and FeCl3-PEO-LiTFSI electrolytes.

Fig. 2

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Balancing the trade-offs between processability (film formation) and electrochemical properties, 3 wt% FeCl3 was determined as the optimal doping ratio (denoted as FeCl3-PEO-LiTFSI). Notably, the optimized FeCl3-PEO-LiTFSI electrolyte exhibited a significantly softer feature. AFM measurements in DMT modulus mode (Fig. 2g) showed that the modulus decreased sharply from 30.8 MPa for PEO-LiTFSI to only 3.1 MPa for FeCl3-PEO-LiTFSI. Meanwhile, depth-sensing nanoindentation revealed that the FeCl3-PEO-LiTFSI film required a maximum load of only 10.1 mN at 300 nm indentation depth, much lower than the 28.6 mN needed for the PEO-LiTFSI film (Fig. 2h). This remarkable soft characteristic is critical for the following mechanism study, as it allows us to rule out the possibility that the improved battery lifespan originates from mechanical inhibition.

2.2. Investigation of Li deposition morphology and its underlying mechanism

The Li morphology was investigated after galvanostatic cycling at different current densities in symmetric cells, which varied significantly between FeCl3-PEO-LiTFSI and PEO-LiTFSI systems (Fig. 3a). The essential distinction between spherical Li growth and conventional dendritic growth lies in the nucleation behavior. In conventional PEO-LiTFSI systems, random defect-induced Li nucleation causes local electric-field amplification and Li+ depletion, leading to anisotropic dendritic growth. In FeCl3-PEO-LiTFSI, uniformly distributed Fe–O/Cl centers provide lithiophilic nucleation sites, where Cl− enriches Li+ and redox-active Fen+ regulates charge transfer. This cation–anion synergy promotes homogeneous Li nucleation and isotropic spherical growth, thereby suppressing dendrite formation. In the FeCl3-PEO-LiTFSI system, Li spheres were consistently observed with particle sizes ranging from 0.2 to 1 µm when cycled at 0.1 mA cm−2. In contrast, the PEO-LiTFSI system demonstrated a significantly smaller size of deposited Li, with sizes around 100 nm. The larger, uniformly distributed Li spheres in the FeCl3-PEO-LiTFSI system indicate a more favorable environment that promotes fewer but larger nuclei, thereby enhancing interfacial stability. After 100 cycles, the morphology of deposited Li was re-assessed. Obviously, the Li morphology in the FeCl3-PEO-LiTFSI system maintained a spherical morphology, with particle size ranging from 5 to 20 µm. The size evolution of Li spheres aligns with classical Ostwald ripening behavior—a fundamental phenomenon describing the thermodynamically driven growth of larger particles at the expense of smaller ones. For deposited Li spheres, smaller particles exhibit a higher surface energy (and thus a higher chemical potential) than larger particles. Upon cycling, small Li spheres tend to dissolve first and redeposit on the surface of larger Li spheres, leading to a gradual increase in the average Li sphere size. As a comparison, the Li morphology in the PEO-LiTFSI systems became disordered and irregular, causing interfacial deterioration. Neither system exhibited the growth of Li dendrites and short circuits, likely due to the short cycling time and low current density.

Fig. 3. (a) Morphologies of deposited Li after cycling with FeCl3-PEO-LiTFSI or PEO-LiTFSI electrolytes: cycled at 0.1 mA cm−2 or 0.3 mA cm−2 for 1 or 100 cycles. (b) Morphologies of deposited Li after cycling with FeCl3-PEO-LiTFSI electrolyte under dynamic current protocols: first 10 cycles at 0.1 mA cm−2, followed by 10 cycles at 0.5 mA cm−2, and finally 10 cycles at 0.1 mA cm−2. (c) Energy-dispersive X-ray spectroscopy (EDS) mapping images of deposited Li at 0.1 mA cm−2 for 100 cycles.

Fig. 3

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It is well known that high current density tends to promote Li dendrite formation by triggering Li+ depletion at the interface. To showcase the advantages of the FeCl3-PEO-LiTFSI system, galvanostatic measurement was performed under a much higher current density (0.3 mA cm−2). As shown in Fig. 3a, in the FeCl3-PEO-LiTFSI system, the size of deposited Li spheres was around 100–150 nm after the initial cycle, which increased to 0.2–1 µm after 100 cycles. The size of Li spheres deposited at 0.3 mA cm−2 is much smaller than that deposited at 0.1 mA cm−2. In accordance with Butler–Volmer equation and classical nucleation theory, a higher current density gives rise to a larger overpotential and a stronger nucleation driving force. This enhanced driving force promotes more frequent nucleation events and reduces the critical nucleation radius, resulting in the formation of smaller Li spheres at 0.3 mA cm−2. Consistent with this trend, further increasing the current density to 0.5 mA cm−2 resulted in even smaller Li spheres with a radius of ∼80 nm (Fig. 3b). Interestingly, when the current density was switched back from 0.5 to 0.1 mA cm−2, the radii of the Li deposits increased to 0.2–1 µm, mostly recovering to the size range observed under the original 0.1 mA cm−2.

Different from the stable performance of the FeCl3-PEO-LiTFSI system, the polarization overpotential of the PEO-LiTFSI system displayed fluctuations and soft short-circuit upon cycling at 0.3 mA cm−2 (Fig. S9). After the initial cycle, the Li morphology exhibited an irregular and wrinkled shape, interspersed with some sharp-edged structures (Fig. 3a). After 100 cycles, prominent Li dendrites with sharp tips and elongated tubular morphology were observed in the PEO-LiTFSI system (Fig. 3a). The comparative results proved the effectiveness of FeCl3-PEO-LiTFSI in achieving large-size spherical Li deposits and consequently a stable interface. Energy-dispersive X-ray spectroscopy (EDS) mapping was performed to analyze the spatial elemental distribution (Fig. 3c). The results showed that the surface of Li spheres was rich in Fe, Cl, F, and O elements. The enrichment of Fe, Cl, and O elements revealed that Li+ ions were preferentially reduced around Fe–O/Cl centers at the interface, suggesting their critical role in regulating Li deposition. The signals related to F and O elements originated from inorganic compounds such as LiF, Li2O, and Li2CO3 within the SEI layer, which will be discussed in detail in the following section. Overall, the persistent spherical morphology across all test conditions and the direct correlation with spatial distribution confirm the efficacy of Fe–O/Cl centers in facilitating the formation of spherical Li deposits. In summary, the Fe–O/Cl centers regulate Li deposition from the nucleation stage through chemical modulation rather than relying on mechanical blocking. This regulation converts Li growth from defect-dominated anisotropic deposition to site-regulated isotropic spherical deposition, thereby enabling effective dendrite suppression even at a low modulus of 3.1 MPa.

Nucleation overpotential (ηnuc.) and exchange current density (j0) were key descriptors for the Li deposition process. The ηnuc. for the Li|PEO-LiTFSI|Cu cell exceeded 20 mV, while for the Li|FeCl3-PEO-LiTFSI|Cu cell, it was nearly zero (Fig. 4a). The differences arise from the electrolyte-induced interface reconstruction. In traditional PEO electrolytes, the space-charge effect tends to induce interfacial Li+ depletion, increase mass-transport polarization, and raise the Li nucleation overpotential.24,36 In the FeCl3-PEO-LiTFSI system, the coordinated Cl− species in Fe–O/Cl centers act as local Li+-enriching sites, which counterbalance space-charge-induced Li+ depletion and homogenize the interfacial Li+ concentration and flux.37 Together with the regulated charge-transfer process mediated by redox-active Fen+ centers, this interfacial Li+ enrichment effect lowers the mass-transport barrier for Li reduction, thus contributing to the ultra-low nucleation overpotential. Based on the classical equations for heterogeneous nucleation, the critical radius of Li nuclei (rcrit) can be expressed as: rcrit = 2γVm/(F|ηnuc.|), where γ is the surface energy and Vm is the Li molar volume.38,39 This equation indicates an inverse relationship between Li nuclei size and overpotential. Therefore, the significantly lower ηnuc. suggested that larger Li nuclei were formed more easily at the Li|FeCl3-PEO-LiTFSI interface during initial plating, consistent with SEM observation. j0 is another key parameter representing the magnitude of equal anodic and cathodic current densities at the equilibrium potential. From Tafel analysis (Fig. 4b), the j0 of the FeCl3-PEO-LiTFSI system was 0.048 mA cm−2, much lower than that of the PEO-LiTFSI system (0.201 mA cm−2). According to a phase-field model, a lower j0 facilitates the formation of a larger Li nucleus.27,28 Moreover, the equilibrium voltage shifted from −64 mV (PEO-LiTFSI) to −20 mV (FeCl3-PEO-LiTFSI), suggesting a modulated redox kinetics due to interfacial Fe–O/Cl centers.

Fig. 4. (a) Nucleation overpotentials evaluated in Li‖Cu cells at 0.1 mA cm−2. (b) Exchange current densities determined from Tafel plots (scanning rate: 0.5 mV s−1). (c) Charge density difference calculated with n additional electrons (n = 0, 1, or 2) being injected into the whole system. (d) Calculated structures and relative energy with varying Li–Li distances. (e) Critical current densities of FeCl3-PEO-LiTFSI. (f) Li morphologies of cycled batteries using FeBr3-PEO-LiTFSI or FeF3-PEO-LiTFSI electrolytes, respectively.

Fig. 4

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Theoretical calculations were conducted to explore the effects of interfacial Fe–O/Cl centers on Li nucleation and growth (Fig. 4c, d and S10–S14). According to the EXAFS and simulations results, electron-deficient Fen+ coordinated with both ether oxygen from the PEO matrix and adjacent Cl− ions, thereby forming Fe–O/Cl centers. The Fe–O/Cl centers contain redox-active Fen+ and the nucleophilic Cl− with strong Lewis basicity, which collaboratively affect Li+ migration, reduction, and growth. Notably, Fe–O/Cl centers form a multi-atom structure that has a significantly higher size than Li+. When serving as nucleation sites, Fe–O/Cl centers can provide more sites for accommodating more Li+, thereby improving the possibility of growing large Li nuclei. During the initial Li plating process, Li+ migrate and are reduced near the nucleation sites, alongside the electron transfer process. The coordinated Cl− have high electron density and polarizability, which enhanced the Li+ affinity of Fe–O/Cl centers. These Cl− will attract Li+ influx into Fe–O/Cl centers, functioning like an “Li+ pump” to maintain a continuous Li+ supply (Scheme 1). Thus, the local Li+ depletion phenomenon and the associated “tip effect” are effectively suppressed. Next is the Li+ reduction process, wherein the charge transfer process plays an important role. For Fe–O/Cl centers, the valence-variable Fen+ exhibits electron donation/acceptance capability—an attribute that effectively facilitates the electron transfer process. DFT calculations revealed that the Mulliken charge of Fen+ was +2.66 (Fig. 4c). When one or two electrons were introduced into the system, the Mulliken charge decreased to +2.42 and +1.56, respectively. The progressive decrease in the Mulliken charge suggested that electrons were effectively localized around the Fen+ centers, facilitating charge transfer processes and further promoting Li+ reduction. This highlights the crucial role of valence-variable Fen+ as “electron reservoirs” that can store and release electrons in response to the Li+ reduction reaction. Overall, the nucleophilic Cl− act as a “Li+ pump” enriching Li+ to Fe–O/Cl centers, while the redox-active Fen+ serve as an “electron reservoir” facilitating the Li+ reduction process. The interplay between coordinated Cl− and Fen+ enables the initial uniform Li deposition around Fe–O/Cl centers.

The dispersion of reduced Li atoms (denoted as *Li) is critical for determining Li morphology. It is essential to figure out how *Li prefer to behave around Fe–O/Cl centers—whether they gather into metallic aggregates or stay spread out as discrete atoms—during the initial nucleation. To investigate this, two *Li were introduced into the model, and the relative energies were calculated at various distances (Fig. 4d). This analysis aimed to identify the most stable position of *Li, providing insights into their preferred behavior (gathering or spreading). The initial state, where the two *Li are far apart without interaction, was defined as the zero-energy state. With the movement of *Li toward the Fe–O/Cl center, the energy profiles decreased, suggesting that the migration of *Li to Fe–O/Cl centers was thermodynamically favored. The lowest energy (−1.451 eV) was observed when the distance of two *Li was reduced to 3.61 Å, representing the most stable configuration. Notably, the typical length of a Li–Li metallic bond was 3.03 Å. If two discrete *Li were artificially brought closer to form Li–Li bonds, the relative energy will increase to −1.141 eV. This indicates that *Li tend to stay isolated around Fe–O/Cl centers at the very initial reduction stage, rather than aggregating into metallic Li. As long as the surroundings of Fe–O/Cl centers are not fully occupied, *Li will prefer to stay spread out as isolated atoms. Given that the size of Fe–O/Cl centers is much larger than that of individual *Li, the initially dispersed *Li will form a layer surrounding the Fe–O/Cl centers. This state eventually evolves into metallic Li once more *Li are introduced. Importantly, the early-stage preference for forming dispersed *Li around Fe–O/Cl centers is the primary reason for the growth of spherical Li deposits and the mitigation of Li dendrite formation. This mechanism is further validated by the significant enhancement in the critical current density (CCD) (Fig. 4e), a key parameter for quantifying the ability to suppress Li dendrites. Symmetrical cells utilizing FeCl3-PEO-LiTFSI electrolyte achieved a high CCD of 2.8 mA cm−2, significantly higher than that using PEO-LiTFSI electrolyte (0.6 mA cm−2). The improved CCD is a direct result of optimized Li nucleation and growth behavior (the regulated spherical nucleation enabled by Fe–O/Cl centers), which ensures more uniform Li deposition and thereby reduces the risk of dendrite-induced short circuits.

It should be noted that, according to depth-sensing nanoindentation and AFM-DMT results (Fig. 2f and g), the FeCl3-PEO-LiTFSI electrolyte exhibits significantly reduced mechanical strength. This observation further confirms that the enhanced Li dendrite suppression originates from the regulated spherical Li deposits mediated by Fe–O/Cl centers, rather than the physical inhibition effect reported in previous literature. A comparison of the modulus/strength and CCD values of recently reported PEO-based solid polymer electrolytes is summarized in Table S4, further highlighting that FeCl3-PEO-LiTFSI achieves a high CCD despite its low modulus. Other ferric halides (FeBr3 and FeF3) were selected as additives for further study, producing FeBr3-PEO-LiTFSI and FeF3-PEO-LiTFSI electrolytes. After 100 galvanostatic cycles at 0.1 mA cm−2, spherical Li deposits were observed across all halide-doped systems (Fig. 4f)—a phenomenon that aligns with the aforementioned Li nucleation mechanism. Notably, the size of these Li deposits exhibited distinct variations, with the radii of the Li spheres showing a clear anion-dependent trend: FeCl3-PEO-LiTFSI (∼0.2 to 1 µm) > FeF3-PEO-LiTFSI (∼200 to 300 nm) > FeBr3-PEO-LiTFSI (∼100 nm). This size discrepancy is hypothesized to arise from the different nucleophilicity of halide anions and structural differences in Fe–O/X (X represents F/Cl/Br) centers. Further investigations are underway in our group to elucidate the detailed underlying mechanisms. The above experimental and theoretical results confirm that Fe–O/Cl centers are conducive to regulating Li plating/stripping behaviors. The collaborative effects of the Li+-affinity Cl− and redox-active Fen+ and the favourable energy dynamics for high *Li dispersion significantly promote the formation of spherical Li deposits and enable an ultra-stable interface.

2.3. Chemistry of interfacial properties

The chemical composition and structure of the SEI layer formed on the Li surface are crucial in influencing Li nucleation and interfacial diffusion. Depth-profiling XPS measurements were performed to investigate the interfacial chemical environment (Fig. 5a–e, S15 and S16). Compared with the Li|PEO-LiTFSI interface, the SEI layer in the Li|FeCl3-PEO-LiTFSI interface contained a significantly higher proportion of inorganic compounds, specifically, LiF, Li2O, and Li2CO3. As illustrated in Fig. 5b, the O 1s spectrum can be deconvoluted into three peaks at 533.5, 531.2, and 528.3 eV, corresponding to organic –C–O– and inorganic Li2CO3, and Li2O, respectively. Notably, the Li2O signal was obviously observed at the Li|FeCl3-PEO-LiTFSI interface and was completely absent at the Li|PEO-LiTFSI interface (Fig. 5a and b). Depth-profiling XPS further revealed a high Li2O content throughout the entire SEI layer at the Li|FeCl3-PEO-LiTFSI interface, which reduced the Li+ diffusion energy barrier and promoting lateral Li deposition.30 The F 1s XPS spectra (Fig. 5c and d) displayed two peaks at 689.0 and 685.0 eV, corresponding to –CF3 and LiF, respectively. At the Li|FeCl3-PEO-LiTFSI interface, LiF demonstrated a consistent and high content throughout the SEI layer, unlike the uneven distribution observed in the Li|PEO-LiTFSI interface. The enrichment and hybridization of LiF and Li2O within the SEI layer are critical for enhancing surface ion migration and promoting lateral Li growth.30 These structural features benefit the overall uniformity of Li deposition. Given that FeCl3-PEO-LiTFSI may react to form a Fe-containing fragment in SEI, XPS was performed to probe the existing states of Fen+ (Fig. 5e). As the sputtering time increased, the Fe 2p signal became more pronounced, indicating its gradient distribution. It was observed that Fen+ demonstrated multiple valence states, likely due to its unique electronic configuration and its capability to exist with different oxidation states. A small proportion of metallic Fe was detected, attributed to the strong reducing ability of Li metal. The presence of Fe-containing complexes with variable valence states can act as catalysts during the initial stage of SEI formation, which was analyzed in the following section.

Fig. 5. Interfacial chemical environment analysis. (a–d) O 1s and F 1s spectra of FeCl3-PEO-LiTFSI|Li and PEO-LiTFSI|Li interfaces, respectively. (e) Fe 2p spectra of the FeCl3-PEO-LiTFSI|Li interface. (f) TOF-SIMS ion surface mapping and depth profile corresponding to the 3D reconstruction distribution of different components in the FeCl3-PEO-LiTFSI|Li interface after cycling at 0.1 mA cm−2 for 100 cycles. (g and h) Energy profile associated with LiTFSI decomposition towards LiF and Li2O formation, respectively.

Fig. 5

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Time-of-flight secondary ion mass spectrometry (TOF-SIMS) was further conducted to confirm the three-dimensional chemical distribution of the SEI layer. The results showed the signals of Fe+, Cl−, LiF2−, and Li2O+, reflecting the existence of Fe–O/Cl, LiF, and Li2O components inside the interfacial layer. The relative abundance of these species followed the order: LiF2− > Cl− > Li2O+ > Fe+ (Fig. 5f). Upon sputtering, distinct changes in the depth profiles were observed: the intensities of LiF2− and Cl− increased significantly with depth. The higher LiF content in the inner layers was beneficial for passivating the electrode and suppressing side reactions. In contrast, the intensities of Li2O+ and Fe+ remained relatively stable, suggesting a more uniform distribution of these species throughout the SEI layers. The TOF-SIMS results, in conjunction with XPS, confirm the hybrid nature of SEI formed at the Li|FeCl3-PEO-LiTFSI interface. Overall, the inner layers are characterized by the enrichment of LiF, ensuring robust electrode passivation, while the uniform distribution of Li2O throughout the SEI enhances ionic transport and supports uniform Li deposition.30 Since LiF and Li2O are decomposition products of LiTFSI, the stronger signals for LiF and Li2O indicate an accelerated decomposition of LiTFSI at the Li|FeCl3-PEO-LiTFSI interface.

DFT calculations were performed to investigate the catalytic effect of Fen+ centers on LiTFSI decomposition and LiF/Li2O formation. Three coordination configurations between Fen+ and LiTFSI were identified, namely η2(O, O), η2(O, N), and η3(O, O, N), with binding energies of −197.53 eV, −196.08 eV, and −195.98 eV, respectively (Fig. S11). Among these, the η2(O, O) was determined to be most thermodynamically favorable and thus selected to explore LiTFSI decomposition pathways. Under the catalytic effect of Fen+, LiTFSI decomposition proceeded via two distinct pathways (Fig. 5g and h). For the LiF formation pathway, the clean surface (state 1, labelled as *) initially provided active sites for LiTFSI adsorption, adopting the η2(O, O) configuration (state 2). This adsorption step lowered the system's energy from 0 eV (state 1) to −8.06 eV (state 2), confirming its thermodynamically favorability. Subsequent cleavage of LiTFSI and desorption of LiOSNSO resulted in the F3C*-Fe-*CF3 intermediate (state 3), further decreasing the energy to −10.82 eV. As the reaction progressed, the system transitioned to F2C*-Fe-*CF2 (state 4) at −19.25 eV, where the first LiF was generated. The reaction continued with additional LiF formation (state 5), reducing the energy to −28.12 eV, and then reached the final state (2*CO) with an energy of −39.40 eV, completing the LiF formation. For the Li2O formation pathway, a similar reaction sequence began with LiTFSI dissociation in the η2(O, O) configuration (state 2) and F3C*-Fe-*CF3 intermediate formation (state 3). Subsequent release of LiCF3 transitioned the system to the F3C*-Fe-*O state (state 4), lowering the energy to −11.13 eV, marking the beginning of oxygen-related intermediate steps. A second LiCF3 release led to the formation of the *O–O* state (state 5), with a sharp energy drop to −16.58 eV, priming oxygen atoms for Li2O formation. The process concluded in the *O state (state 6), where Li2O was fully formed and stabilized on the surface, achieving the lowest energy of −26.09 eV. Detailed Gibbs free energy changes for each step of the above reactions are provided in Fig. S12 and S13. For comparison, control simulations without Fen+ centers showed that LiF formation faced a substantial energy barrier (ΔG = 5.13 eV) at the third intermediate step, while Li2O formation did not occur at all (Fig. S14). These theoretical results were consistent with those from XRD and TOF-SIMS.

The catalytic role of Fen+ sites was further validated through systematic experiments. First, symmetric Li cells cycled at 0.1 mA cm−2 for 1 cycle were disassembled for interfacial characterization. As shown in Fig. S15, distinct peaks corresponding to LiF and Li2O were detected after the initial cycle in depth-profiling XPS spectra, confirming the effectiveness of Fen+ in accelerating an LiF/Li2O-rich interfacial layer. Additionally, the Fe 2p spectrum after the initial cycle (Fig. S15) was similar to that after 100 cycles (Fig. 5e), suggesting that Fen+ centers were relatively stable without continuous reduction during cycling. After a slight initial reduction, Fen+ maintained its chemical states without further reduction, which was quite important. To further clarify the influence of FeCl3 loadings, control experiments were conducted with FeCl3 content reduced from 3 wt% to 1 wt%. Symmetric cells employing FeCl3(3%)-PEO-LiTFSI or FeCl3(1%)-PEO-LiTFSI electrolytes were disassembled after 100 cycles at 0.1 mA cm−2. As shown in Fig. S16, compared with the additive-free PEO-LiTFSI electrolytes, the contents of LiF and Li2O in the Li surface increased progressively with the increase in FeCl3 content. This result confirms that a higher Fen+ concentration enhances the catalytic decomposition of TFSI− into LiF and Li2O, consistent with the proposed catalytic mechanism. Overall, the above findings underscore the critical role of Fen+ in reducing energy barriers for uniform Li deposition and enabling the formation of a LiF/Li2O-rich SEI layer that promotes uniform Li plating/stripping.

2.4. Battery performances

To evaluate the tolerance of Li dendrites, Li‖Li symmetric cells were assembled and cycled under various current densities. The Li|FeCl3-PEO-LiTFSI|Li cell demonstrated stable cycling at current densities of 0.1, 0.2, 0.3, 0.4, and 0.5 mA cm−2, with corresponding polarization values of 66, 132, 168, 220, and 275 mV, respectively (Fig. 6a). These consistently low polarization values indicate the enhanced ionic conductivity and interfacial stability of the FeCl3-PEO-LiTFSI electrolyte. In contrast, the symmetric cell with PEO-LiTFSI electrolyte exhibited significantly poorer performance due to short-circuiting after approximately 60 cycles at 0.3 mA cm−2. The improved performance in the FeCl3-PEO-LiTFSI electrolyte system can be attributed to its regulated spherical Li morphologies and SEI structures as discussed above. Notably, the Li|FeCl3-PEO-LiTFSI|Li cell maintained stable cycling performance for over 3000 hours at 0.1 mA cm−2 (Fig. 6b). This exceptional longevity demonstrated the remarkable dendrite suppression capabilities of the FeCl3-PEO-LiTFSI electrolyte. In contrast, the PEO-LiTFSI-based symmetric cell experienced a short circuit after only 164 h due to the dendrite-induced short circuits. Post-cycling EIS curves were measured to monitor the dynamic impedance evolution during cycling (Fig. S17). For the PEO-LiTFSI system, the symmetric cell initially exhibited a low resistance of 58 Ω. After 100 cycles, the cell resistance increased to 75 Ω, reflecting the interfacial deterioration caused by uneven Li deposition. After 500 cycles, the resistance decreased to 52 Ω, indicating the occurrence of micro-shorts, consistent with the dropped polarization voltage observed in the galvanostatic cycling curves. In contrast, the Li|FeCl3-PEO-LiTFSI|Li cell displayed ultra-stable EIS profiles with a consistent resistance of ∼150 Ω throughout the entire cycles. The higher initial cell resistance of the Li|FeCl3-PEO-LiTFSI|Li cell can be ascribed to the in situ formed interfacial layers. Importantly, the absence of resistance fluctuation and ultra-long lifespan verified the effectiveness of FeCl3-PEO-LiTFSI in achieving an ultra-stable interface.

Fig. 6. Performances of the symmetrical and full cells. (a) Long-term cycling stabilities of Li|FeCl3-PEO-LiTFSI|Li and Li|PEO-LiTFSI∣Li symmetric cells at different current densities. (b) Ultralong cycling performance of the Li|FeCl3-PEO-LiTFSI|Li and Li|PEO-LiTFSI|Li symmetric cells at 0.1 mA cm−2 and 60 °C. (c) Cycling performance of Li|FeCl3-PEO-LiTFSI|LiFePO4 and Li|PEO-LiTFSI|LiFePO4 at 0.5C and 60 °C. (d) Comparison of the battery performance of this work with previous studies.

Fig. 6

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To evaluate the practical feasibility, all-solid-state batteries were assembled using LiFePO4 (LFP) as the cathode and Li metal as the anode. The Li|FeCl3-PEO-LiTFSI|LFP cell delivered an impressive initial specific capacity of 138.4 mAh g−1 at 0.5C and maintained 129.4 mAh g−1 after 1000 cycles (Fig. 6c), demonstrating exceptional long-term cycling stability. Compared to recently reported results in the literature, this performance was outstanding (Fig. 6d).40–47 Furthermore, the corresponding voltage profiles of the Li|FeCl3-PEO-LiTFSI|LFP cell showed negligible differences between the 1st, 100th, 200th, and 1000th cycles, indicating highly stable charge–discharge behavior (Fig. S18). In contrast, the Li|PEO-LiTFSI|LFP cell exhibited poor cycling stability, with its coulombic efficiency declining rapidly after 200 cycles. Additionally, the corresponding charge–discharge profiles became increasingly unstable, highlighting significant performance degradation. These results underscore the critical role of FeCl3-modified electrolyte in suppressing dendrite formation and enhancing the stability and longevity of all-solid-state batteries.

3. Conclusions

The incorporation of FeCl3 into PEO-LiTFSI significantly enhances the electrochemical performance through the formation of Fe–O/Cl centers. The redox-active Fen+ act as electron reservoirs that optimize the electron transfer process, while the Li+-affinity Cl− function as Li+ pumps that enable sufficient Li+ supply. During Li electroplating, the reduced *Li prefers to stay spread out as isolated atoms surrounding the large-sized Fe–O/Cl centers, facilitating the growth of Li spheres. Furthermore, Fen+ facilitated LiTFSI decomposition to construct a hybrid Li2O/LiF-rich SEI layer at the Li/electrolyte interface. Collaboratively, FeCl3-PEO-LiTFSI electrolyte achieves an ultra-stable interface that enables a high critical current density of 2.8 mA cm−2, long-term cycling stability of 3000 hours in Li symmetric cell, and a capacity retention of 93.5% after 1000 cycles in LiFePO4-based full cells. This work demonstrates the significant potential of FeCl3-PEO-LiTFSI electrolyte as a promising candidate for advanced Li metal battery applications. The elucidated mechanism will pave a new way for regulating Li deposition behavior and designing high-performance polymer electrolytes.

Author contributions

T. Z. and Y. X. designed and performed the experiment and wrote the original manuscript. Z. H. performed the experiment. J. Z. and L. G. participated in the discussion and revised the manuscript. R. Z. designed and funded the experiments, analyzed the results, and revised the manuscript. C. Y. analyzed part of the data. S. H. and Q. X. supported the experiments and revised the manuscript.

Conflicts of interest

The authors declare no conflicts of interest.

Supplementary Material

SC-OLF-D6SC04034A-s001

Acknowledgments

This research was supported by the Guangdong Basic and Applied Basic Research Foundation (2025A1515010309), Guangdong Grants (2021ZT09C064), and National Natural Science Foundation of China (42202137). The authors acknowledge the support from the Major Science and Technology Infrastructure Project of Material Genome Big-science Facilities Platform supported by the Municipal Development and Reform Commission of Shenzhen, Guangdong Provincial Key Laboratory of Energy Materials for Electric Power, and Shenzhen Key Laboratory of Micro/Nanoporous Functional Materials (SKLPM) (ZDSYS20210709112802010), BL16U1 and BL17B1 beamlines of the Shanghai Synchrotron Radiation Facility, and the Open Research Fund of Guangdong Advanced Carbon Materials Co., Ltd (Kargen2024B0904).

Data availability

Datasets are available from the corresponding author upon reasonable request. The data supporting this article have been included as part of the supplementary information (SI). Supplementary information: experimental and computational details, supplementary characterization and electrochemical data, and supporting figures and tables. See DOI: https://doi.org/10.1039/d6sc04034a.

Notes and references

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Associated Data

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Supplementary Materials

SC-OLF-D6SC04034A-s001
SC-OLF-D6SC04034A-s001.pdf (2MB, pdf)

Data Availability Statement

Datasets are available from the corresponding author upon reasonable request. The data supporting this article have been included as part of the supplementary information (SI). Supplementary information: experimental and computational details, supplementary characterization and electrochemical data, and supporting figures and tables. See DOI: https://doi.org/10.1039/d6sc04034a.

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