A MXene Colloid Electrolyte With Spatiotemporally Continuous Ion Transport for Fast‐Charging Lithium Metal Batteries

ABSTRACT

Dendrite growth and interfacial instability in lithium metal batteries are closely associated with imbalanced lithium‐ion transport, particularly at high current densities. Such an imbalance reflects the loss of spatiotemporal continuity in Li⁺ supply during cycling, manifested by localized amplification of potential and insufficient Li⁺ replenishment. Herein, a sulfonated MXene (Ti₃C₂(SO₃H)_x) colloid electrolyte is developed to regulate Li⁺ transport through the concurrent modulation of solvation structure and electric potential distribution. The ─SO₃H surface terminations alter the Li⁺ coordination environment and lower the interfacial desolvation barrier, while conductive MXene sheets redistribute local potential and current density, jointly shaping Li⁺ transport behavior. This electrolyte mitigates spatial localization of ion flux and sustains interfacial Li⁺ replenishment under high current densities, leading to uniform Li metal deposition. As a result, the MXene colloid electrolyte achieves a high Li⁺ transference number of 0.89, facilitating efficient Li⁺ transport. Consequently, Li||NCM622 cells employing the MXene colloid electrolyte achieve outstanding fast‐charging capability and cycling durability at 10 C, delivering twice the cycle life of cells with the baseline electrolyte under equivalent capacity retention. This study shows that the MXene colloid electrolyte design achieves spatiotemporally continuous Li⁺ transport, providing a guiding framework for designing fast‐charging lithium metal battery electrolytes.

Spatiotemporally continuous Li⁺ transport is established in lithium metal batteries using a sulfonated MXene colloid electrolyte. Surface ─SO₃H groups promote rapid Li⁺ desolvation, sustaining timely interfacial ion supply under fast charging, while conductive MXene nanosheets moderate local potential gradients to maintain spatially even ion distribution. This continuity supports uniform lithium metal deposition and stable 10 C cycling in Li||NCM622 cells.

1. Introduction

The rapid growth of portable electronics, electric vehicles, and grid‐scale energy storage has placed increasing demands on rechargeable batteries with substantially higher energy density and improved reliability. Conventional lithium‐ion batteries (LIBs), though technologically mature, are approaching their theoretical energy density limits under current material systems and cell architectures [1, 2, 3]. Lithium metal (Li⁰), with an exceptionally high theoretical specific capacity (3860 mAh g⁻¹), low density (0.534 g cm⁻³), and the most negative potential (−3.04 V vs the standard hydrogen electrode), makes lithium metal batteries (LMBs) capable of achieving energy densities of up to ∼400 Wh kg⁻¹ [4, 5]. Nevertheless, practical application of LMBs remains severely hindered by uncontrolled dendrite growth and parasitic side reactions, which progressively reduce coulombic efficiency, shorten cycling life, and even raise serious safety concerns [6, 7, 8]. These issues fundamentally originate from sluggish Li⁺ transport kinetics across the solid–liquid interface, coupled with the intrinsic instability of lithium metal. Under fast‐charging conditions with highly accelerated ion transport demands, these failure phenomena of LMBs become more pronounced.

In‐depth analysis of the Li⁺ transport mechanisms reveals that macroscopic degradation behaviors such as dendrite growth, capacity decay, and even short circuits stem from imbalanced Li⁺ transport at the microscopic scale [9, 10, 11]. From a transport perspective, Li⁺ migration can be described by the Nernst–Planck equation, in which the ion flux is governed by two driving forces, diffusion and electromigration [12, 13]. Importantly, in working batteries, the critical challenge lies not in the absolute values of these driving forces, but in their spatiotemporal inhomogeneity [14, 15]. Specifically, at high charging rates, diffusion and electromigration driving forces are amplified to meet the large current densities, making Li⁺ transport increasingly sensitive to their spatial and temporal continuity. However, rapid Li⁺ consumption at the solid–liquid interface leads to local depletion during cycling. When the Li⁺ transport becomes uneven in space or cannot be sustained continuously in time, localized concentration polarization arises, resulting in larger concentration overpotentials and intensified local electric fields. During cycling, the spatiotemporally inhomogeneous distribution of Li⁺ amplifies local ion flux at protrusions on the electrode surface, thereby promoting the formation of lithium dendrites [16, 17]. Notably, the extent to which spatiotemporally discontinuous ion transport is strongly influenced by the local solvation environment, which governs Li⁺ coordination and desolvation kinetics. Under fast‐charging conditions, Li⁺ depletion at the solid–liquid interface is significantly aggravated, thereby accelerating performance degradation and cell failure. Therefore, achieving spatiotemporally continuous ion transport is considered a central challenge in enabling fast‐charging LMBs.

Despite extensive studies on stabilizing Li⁰ anodes, most strategies fail to comprehensively enhance the spatiotemporal continuity of Li⁺ transport, thereby limiting long‐term stability under practical conditions. For instance, in electrolyte optimization, employing high‐concentration electrolytes and designing anion‐coordinated solvation structures can enhance interfacial stability [18, 19, 20, 21]. However, these strategies are often accompanied by sharply increased viscosity, elevated cost, and even reduced bulk ionic conductivity, which are unfavorable for sustained and rapidly responsive Li⁺ transport under high current densities. In interfacial engineering, artificial solid‐electrolyte interphase (SEI) or functional interlayers have been developed to promote uniform Li⁺ transport and to physically suppress dendrite growth [22, 23, 24, 25]. Recent studies have shown that incorporating functional materials such as MXene into interfacial layers can enhance lithiophilicity and regulate Li⁺ deposition behavior [26, 27]. Yet, these interfacial engineering methods primarily regulate Li⁺ transport within a confined interfacial region, while ion transport uniformity in the bulk electrolyte remains largely uncontrolled. Consequently, nonuniform Li⁺ distribution gradually develops during prolonged cycling, undermining their effectiveness over extended operation at high‐rate conditions. In electric‐field engineering, designing 3D current collectors with tailored geometries has been explored to redistribute the electric field and guide Li⁺ toward recessed regions [28, 29, 30]. Nonetheless, these strategies based on 3D current collectors often operate at fixed geometric length scales defined by the electrode architectures and are rarely integrated with solvation regulation or interfacial ion‐transport kinetics, making it challenging to maintain spatiotemporal continuity of Li⁺ transport across the solid–liquid interface during long‐term cycling.

Electric‐field regulation represents a conceptually sound approach for directing Li⁺ transport and deposition behavior. Since the electric field acts as the driving force for Li⁺ electromigration, locally intensified fields can drive preferential Li⁺ migration toward specific regions, resulting in uneven ion flux and nonuniform deposition. A smoother potential gradient, by contrast, favors more homogeneous Li⁺ transport and deposition [31, 32]. However, most existing electric‐field‐based strategies are predominantly confined to electrode architectures or the electric double layer, where their influence is spatially localized and difficult to extend into the electrolyte bulk, hindering the realization of spatiotemporally homogeneous Li⁺ transport throughout the cell. Introducing dispersed functional species into the electrolyte may provide a feasible route to modulate the potential distribution and to influence ion transport over a broader region [28, 33]. Two‐dimensional MXene nanosheets are well suited to this strategy, as their high electronic conductivity enables redistribution of local electric potential, while their surfaces allow the introduction of polar functional groups to regulate Li⁺ coordination. Together, these attributes directly correspond to the spatial and temporal inhomogeneities in Li⁺ transport identified above.

In this work, we develop a sulfonated MXene (Ti₃C₂(SO₃H)_x) colloid electrolyte to realize the spatiotemporally continuous transport strategy (Figure 1). The ─SO₃H surface groups alter the Li⁺ coordination environment, lowering the desolvation barrier, thereby improving interfacial Li⁺ transport kinetics and alleviating local Li⁺ depletion near the electrode. Accordingly, fast Li⁺ transport at the solid–liquid interface meets the timescale requirements for rapid ionic response under high‐rate charging. Meanwhile, the highly conductive MXene sheets undergo field‐induced polarization and redistribute the local potential, helping homogenize the spatial Li⁺ distribution. These two effects endow the MXene colloid electrolyte (MCE) with spatiotemporally continuous ion transport, resulting in faster and more uniform Li⁺ flux. As a result, MCE enables dense and stable lithium deposition and significantly improves the fast‐charging performance of Li||NCM622 cells. This work demonstrates a coordinated transport‐regulation strategy and offers mechanistic insights for designing stable lithium metal batteries under fast‐charging conditions.

FIGURE 1.

Schematic illustration of the sulfonated MXene colloid electrolyte for regulating Li⁺ transport in LMBs. The upper panel depicts the MXene colloid electrolyte (MCE), while the lower panel shows the corresponding behavior in the baseline electrolyte (BE) for comparison. The sulfonated groups provide adsorption sites that alter the Li⁺ coordination environment and lower the interfacial desolvation barrier, enabling fast desolvation and prompt Li⁺ availability. Meanwhile, the conductive MXene nanosheets induce field‐mediated modulation of electric potential gradients and current pathways, leading to an even ion distribution. The coupled regulation suppresses dendrite growth, promoting uniform lithium deposition under high current densities.

2. Results and Discussion

2.1. Structure and Properties of MXene Colloid Electrolyte

Ti₃C₂T_x (T_x = ─O, ─OH, etc.) was synthesized via in situ etching and intercalation (Figure S1). As shown in the X‐ray diffraction (XRD) pattern in Figure 2a, the delaminated Ti₃C₂T_x no longer exhibited characteristic peaks of the Ti₃AlC₂ precursor. After water intercalation, the (002) peak shifted markedly to lower angles, indicating an expanded interlayer spacing and confirming successful exfoliation into few‐ and single‐layered nanosheets.

FIGURE 2.

Structure and colloidal characteristics of sulfonated Ti₃C₂(SO₃H)_x MXene. (a) XRD patterns of Ti₃AlC₂, Ti₃C₂T_x and Ti₃C₂(SO₃H)_x. (b) FT‐IR spectra of Ti₃C₂T_x and Ti₃C₂(SO₃H)_x. (c) TEM image of Ti₃C₂(SO₃H)_x. (d) Digital photograph of MCE. (e) Viscosity comparison of MCE and BE.

Sulfonic acid groups were then grafted onto the surface via a diazotization‐coupling reaction [34, 35]. This surface functionalization introduces bulky and polar functional groups that promote stable colloidal dispersion and enhance ion‐functional group interactions in the electrolytes. After surface functionalization, the characteristic (002) peaks shifted leftward from 6.37° to 5.44°, corresponding to an expansion in interlayer spacing from 13.88 to 16.23 Å according to Bragg's law. Such interlayer expansion reflects the successful incorporation of ─SO₃H groups and is beneficial for suppressing nanosheet restacking, which in turn stabilizes the colloidal dispersion. Fourier transform infrared spectroscopy (FT‐IR) further supports this assignment. As depicted in Figure 2b, two new bands appeared between 1140−1020 cm⁻¹, where the vibration at 1122 cm⁻¹ is assigned to the O─S─O stretching of ─SO₃H groups, and the band at 1034 cm⁻¹ corresponds to S─C (aryl) stretching [36, 37, 38]. The energy‐dispersive spectroscopy (EDS) mappings in Figure S2 further confirm the uniform distribution of Ti, S, and O, reinforcing the conclusion that sulfonate groups are homogeneously grafted. The transmission electron microscopy (TEM, Figure 2c) and the scanning electron microscopy (SEM, Figure S3) reveal the small‐sized, thin, and sheet‐like Ti₃C₂(SO₃H)_x, in sharp contrast to the densely packed MAX and the accordion‐like multi‐layered MXene. The suppressed restacking and highly exposed surface morphology provide abundant accessible surface sites for Li⁺ interactions.

Dispersing Ti₃C₂(SO₃H)_x into the electrolyte at the concentration of 0.05 mg mL⁻¹, as determined from concentration screening (Figure S4), produces a stable colloid system, showing a clear Tyndall effect in Figure 2d. In contrast, the baseline electrolyte (BE) showed no optical scattering at any stage (Figure S5). This comparison confirmed that the observed scattering originated from the well‐dispersed MXene nanosheets rather than from the solvent itself. Rheological measurements supported the stable dispersion. The viscosity increased slightly from 4.17 cP (BE) to 4.40 cP (MCE), as shown in Figure 2e. Despite this modest increase, the viscosity remains sufficiently low to preserve rapid ion transport, indicating that the incorporation of Ti₃C₂(SO₃H)_x does not compromise electrolyte fluidity or bulk Li⁺ mobility under practical current densities.

2.2. Regulation of Li+ Desolvation and Potential Distribution

To clarify the role of Ti₃C₂(SO₃H)_x MXene in regulating Li⁺ transport under fast‐charging conditions, the following discussion focuses on how the MXene enables temporally and spatially continuous ion transport at the interface and across the electrolyte. At the molecular level, density functional theory (DFT) calculations were used to evaluate the interactions of Li⁺ with electrolyte anions, solvents, and MXene surface groups (Figure 3a), offering insight into coordination environments relevant to the desolvation behavior. The results showed that Li⁺ binds more strongly to the ─SO₃H terminal groups on the MXene surface (−2.80 eV) than to ethylene carbonate (EC) solvent molecules (−1.95 eV). This stronger adsorption alters the original solvent‐dominated coordination environment, shifting the solvation structure from solvent‐separated ion pairs (SSIPs) toward contact ion pairs (CIPs) and aggregates (AGGs), which are associated with lower desolvation barriers and more facile desolvation at the solid–liquid interface. As a result, Li⁺ can be released from its solvation shell and participate in interfacial charge transfer on a shorter characteristic timescale, which is essential for sustaining rapid ionic response under high current densities. Raman spectroscopy further corroborates this restructuring of the Li⁺ solvation shell. In the 880–920 cm⁻¹ region, the characteristic C─O stretching peaks of EC appear at 896 cm⁻¹ for free‐EC and 906 cm⁻¹ for EC involved in solvation [31]. As shown in Figure S6, Ti₃C₂(SO₃H)_x increased the fraction of free‐EC from 31% to 36%, demonstrating that more solvent molecules are released from the Li⁺ solvation sheath, thereby pointing to a solvation environment more favorable for interfacial desolvation. Consistent changes were also observed in the ion‐pairing states. The band near 720 cm⁻¹ corresponds to SSIPs, which require high desolvation energy at the solid–liquid interface and tend to form an organic‐rich SEI [39]. By contrast, the bands at ∼735 and ∼740 cm⁻¹ represent CIPs and AGGs, both of which desolvate more readily and thus facilitate faster Li⁺ desolvation at the interface [40]. As shown in Figures 3b and 3c, Ti₃C₂(SO₃H)_x increased the relative proportions of CIPs and AGGs (68.5% vs 56.5%) while reducing SSIPs (31.5% vs 37%), supporting a transition from a tightly bound, solvent‐dominated solvation shell toward a more weakly coordinated configuration. This weakened Li⁺ coordination lowered the interfacial desolvation barrier, enabling rapid desolvation and timely ion supply at the solid–liquid interface under high current densities. DFT calculations also indicate that Ti₃C₂(SO₃H)_x may provide an additional surface‐assisted migration process in the bulk phase. The –SO₃H terminations exhibit moderately stronger interactions with Li⁺ than solvent molecules due to their localized electron density. These interactions enable Li⁺ to access the MXene surface transiently during migration, allowing motion along the 2D layers with reduced energetic fluctuations in migration energy. This surface‐assisted migration process may buffer local energetic perturbations under high current densities, providing a microscopic basis for sustaining continuous Li⁺ transport in the bulk electrolyte.

FIGURE 3.

Regulation of Li⁺ transport through solvation restructuring and potential homogenization. (a) Binding energies of Li⁺ with electrolyte components with optimized geometries. (b) Raman spectra of MCE and BE, and (c) corresponding percentages of SSIPs, CIPs, and AGGs. Simulated electric potential distributions (color map) and current density vectors (streamlines) of (d) MCE and (e) BE under an applied potential difference. (f) Schematic illustration of electric potential and current redistribution induced by conductive MXene nanosheets, where the background color represents electric potential intensity. DRT analysis of the symmetric cells with (g) MCE and (h) BE. (i) The fitting results of resistance. Inset: the activation energy of the Li⁺ desolvation process. (j) Tafel plots and corresponding exchange current densities of each electrolyte. (k) Arrhenius plots of ionic conductivity. (l) Li⁺ transference number of MCE.

Finite element simulations with manually set parameters were conducted to further elucidate how Ti₃C₂(SO₃H)_x moderates electric potential distribution and promotes spatially continuous Li⁺ flux across the electrolyte. Figure 3d,e present the potential and current density distributions of MCE and BE, respectively. The color maps denote the local electric potential, with red corresponding to higher values and blue to lower ones, while the surface arrows depict the local current vectors, indicating the spatial redistribution of ion transport pathways. In the BE shown in Figure 3d, the potential exhibits an approximately linear profile across the bulk electrolyte, and the current density is largely aligned along the field direction, forming roughly parallel streamlines. However, such macroscopic uniformity does not preclude local disruptions in ion transport, particularly near surface protrusions where ion consumption and field focusing become amplified under high current densities during cycling [41, 42]. By contrast, the addition of Ti₃C₂(SO₃H)_x leads to a more gently sloping and spatially homogenized potential distribution. This effect arises from the intrinsic polarization of Ti₃C₂(SO₃H)_x, enabled by its high electronic conductivity and strong polarizability, which generates local fields orienting opposite the applied electric field and thereby suppresses sharp local gradients [28, 29]. Analysis of local current density reveals that ionic current preferentially follows the surface and interlayer regions of MXene particles. Polarization of the Ti₃C₂(SO₃H)_x sheets in response to the applied electric field bends and redistributes the current lines, guiding the flow along the particles and resulting in a more spatially continuous current distribution, rather than being concentrated at the protrusions. In addition, Li⁺ flux along the y‐direction (perpendicular to the electrodes) is further facilitated by migration pathways associated with MXene surface terminals (Figure S7). Together, these results indicate that Ti₃C₂(SO₃H)_x effectively modulates local electric potentials and current distribution, leading to more spatially uniform Li⁺ transport in the bulk electrolyte.

Building on these mechanistic insights, electrochemical measurements were carried out to examine whether the temporal and spatial continuity of Li⁺ transport is manifested in macroscopic transport properties. Electrochemical impedance spectroscopy (EIS, Figure S8) showed that, compared with BE, the solid‐state interfacial resistance (R_SEI) and charge transfer resistance (R_ct) of MCE were substantially reduced, indicating facilitated Li⁺ transport across the SEI and a lowered barrier for interfacial charge‐transfer processes. The reduced impedance is consistent with a more spatially continuous ion‐transport environment, as a consequence of homogenized electric potential and current density distribution in the presence of Ti₃C₂(SO₃H)_x. Beyond resistance magnitude, DRT analysis further revealed a distinct shift of the interfacial relaxation process toward shorter characteristic time constants for MCE. As shown in Figure S9, at 25°C, the dominant relaxation peak associated with the SEI‐related process appeared at τ ≈ 7.52 × 10⁻⁴ s for MCE, markedly lower than that of BE (τ ≈ 1.03 × 10⁻³ s). Moreover, this leftward shift of the relaxation peak persists consistently across all investigated temperatures, with MCE exhibiting systematically smaller τ values than BE (Figure S10, Table S1). Such a shift toward shorter relaxation times is commonly associated with faster interfacial response and reduced kinetic hindrance, indicating that Li⁺ transport across the solid–liquid interface proceeds on a shorter characteristic timescale in MCE [43]. The activation energies extracted from temperature‐dependent resistance spectra revealed that MCE exhibited a lower reaction energy barrier (31.29 kJ mol⁻¹) than BE (41.71 kJ mol⁻¹). This reduction reflects a decreased interfacial energy barrier, likely arising from the combined effects of faster desolvation and spatially homogenized electric potential distribution, which together facilitate rapid Li⁺ transport across the interface. Corroborating these findings, the MCE‐based system displayed a higher exchange current density, increasing from 1.69 to 5.84 × 10⁻⁴ A cm⁻² (Figure 3j), signifying a more facile charge‐transfer process. To further validate these findings, cyclic voltammetry (CV) tests at a constant scan rate were conducted (Figure S11). MCE exhibited a more positive reduction peak at 3.71 V compared with 3.64 V for BE, together with a narrower redox peak separation of 0.28 V relative to 0.34 V, indicating a more reversible lithium plating/stripping behavior with reduced resistance and accelerated interfacial kinetics. Consistent with the improved kinetics observed above, MCE exhibited an approximately 2.3‐fold increase in ionic conductivity compared to the BE (0.30 mS cm⁻¹), reaching 0.69 mS cm⁻¹ (Figure 3k), with temperature‐dependent conductivity and Arrhenius analysis (Figure S12) further revealing a distinctly lower activation energy for MCE in the electrolyte bulk, while the lithium transference number also increased significantly (0.89 vs 0.59, Figure 3l; Figure S13). These combined improvements demonstrate more efficient and preferential Li⁺ migration in the bulk phase, which, together with the enhanced interfacial kinetics, converges to sustain temporally continuous and spatially uniform ion flux under high current densities. Together, solvation restructuring and electric potential homogenization impart Li⁺ transport with sustained temporal continuity and spatial uniformity, enabling Ti₃C₂(SO₃H)_x to support uniform lithium deposition under fast‐charging conditions.

2.3. Interfacial Stability and Lithium Deposition Behavior

By modulating both the solvation structure and the local electric potential distribution, MCE enables a continuously sustained Li⁺ flux to the lithium metal surface, leading to markedly enhanced interfacial stability. Under galvanostatic cycling at a current density of 1 mA cm⁻² and an areal capacity of 1 mAh cm⁻², Li||Li symmetric cells furnished with MCE maintained a low and stable overpotential for more than 600 h (Figure 4a), whereas the cells with BE exhibited a pronounced increase in polarization after ∼500 h, followed by short‐circuit failure at a later stage. Similar stability trends became more pronounced at a higher current density (Figure S14). Such extended stability was indicative of a well‐maintained Li⁺ flux near the surface throughout prolonged cycling, which supported reversible interfacial Li evolution without the accumulation of poorly compacted deposits during repeated plating/stripping. The reversibility of lithium plating/stripping was further examined in Li||Cu half cells. MCE‐based cells showed higher coulombic efficiency than BE‐based cells (Figure S15). As the long‐term stability is often predetermined by the earliest stages of Li nucleation, the initial nucleation kinetics were analyzed [44, 45]. The initial nucleation overpotential in MCE symmetric cells was 40 mV lower than that in BE (Figure 4b), consistent with a more weakly coordinated solvation environment that lowered the desolvation energy barrier and facilitated a more continuous and homogeneous Li⁺ flux toward the interface. SEM morphology further corroborated the regulated Li deposition behavior. After 20 h of cycling, the MXene electrolyte system exhibited a smooth surface free of dendrites, whereas BE led to porous, irregular Li deposits with abundant voids, as shown in Figures 4c,d. This pronounced morphological contrast indicated that lithium grew with Li⁺ flux that was well‐maintained and spatially distributed along the deposition front. Similar trends were observed after extended plating/stripping over 50 cycles (Figure S16), providing additional support for the sustained interfacial stability of MCE across multiple current densities. X‐ray photoelectron spectroscopy (XPS) analysis was further performed on cycled lithium electrodes (Figure S17), and the related peak information is summarized in Table S2. Compared with BE, the MCE electrolyte favors the formation of a LiF‐rich SEI, which is likely related to the solvation structure shifting from SSIPs to more anion‐involved CIPs and AGGs, beneficial for stabilizing the lithium interface.

FIGURE 4.

Interfacial stability and lithium deposition behavior in Li symmetric cells. (a) Cycling stability at 1 mA cm⁻² with an areal capacity of 1 mAh cm⁻². Inset: enlarged details showing short‐circuiting in BE. (b) Nucleation overpotential during initial lithium deposition. Surface morphologies of lithium anodes after 20 h of cycling with (c) MCE and (d) BE. (e) Rate performance of Li symmetric cells. (f, g) Simulated electrolyte potential and current density distributions. (f) Electrolyte potential (color map) and current density vectors (streamlines). (g) Electrolyte current density magnitude (color map). The hollowed region represents the initially formed lithium dendrites. The surface color denotes the electrolyte potential in Figure 4f, with warmer colors corresponding to higher potentials, and the arrows on the surface indicate the electrolyte current density vectors. The surface color represents the magnitude of the electrolyte current density in Figure 4g, where red regions correspond to higher local current densities. (h) Current density tolerance test, where BE cells are short‐circuited at 10 mA cm⁻².

The benefits of MCE became more apparent with increasing current density. Cycling stability tests were conducted at different current densities (Figure 4e). Cells based on MCE consistently exhibited smaller and more stable overpotentials, reflecting a stable Li⁺ transport state under elevated current demand. To visualize how Ti₃C₂(SO₃H)_x regulates the spatial distribution and dispersion of ion transport near the lithium interface, COMSOL simulations were performed (Figure 4f–g). In BE, pronounced local potential differences and concentrated current density drove ion flux toward protruding sites, resulting in nonuniform lithium deposition. By contrast, the incorporation of Ti₃C₂(SO₃H)_x reorganized the ion flux distribution near the lithium surface. Its conductive layered structure modulated the bulk potential distribution, which in turn homogenized the interfacial potential and local current density, leading to more uniform lithium deposition. The effectiveness of ion‐transport regulation becomes even more evident under extreme operating conditions. The BE‐based cells exhibited short‐circuiting during the first cycle at 10 mA cm⁻² and 10 mAh cm⁻² due to runaway dendrite growth, whereas MCE enabled stable cycling at the same condition (Figure 4h). This high‐current stability indicates that the MXene electrolyte maintains reliable ion transport and deposition control even under extreme current densities. Taken together, the symmetric cell tests demonstrate that MCE delivers a more regulated Li⁺ flux, as evidenced by lower polarization, reduced nucleation overpotential, smoother Li deposits, and stable operation under high‐current densities. These features are consistent with the spatiotemporally continuous Li⁺ transport enabled by MCE and lead to a stabilized anode surface that directly supports the high‐rate operation in lithium metal batteries discussed in the following section.

2.4. Electrochemical Performance of Fast‐charging Lithium Metal Batteries

To further assess how the stabilized lithium behavior translates to practical lithium metal batteries, cycling tests were conducted in cells paired with LiFePO₄ (LFP) and LiNi_0.6Co_0.2Mn_0.2O₂ (NCM622) cathodes. When cycled at 2 C, the MCE‐based Li||LFP cells exhibited exceptional cycling stability, retaining 90% of their initial capacity after 3000 cycles (Figure 5a). In contrast, the BE‐based Li||LFP cells suffered significant capacity fading after 1000 cycles, demonstrating the substantial improvement in long‐term stability enabled by MCE. The charge/discharge profiles in Figure 5b and Figure S18 clearly depict the specific capacity across different cycle numbers. The discharge specific capacity of MCE‐based cells decreased from 146 to 132 mAh g⁻¹, with an average per‐cycle capacity decay of 0.003%, whereas BE cells exhibited a far higher fading, dropping from 134 to 108 mAh g⁻¹ over 2000 cycles. Rate performance curves (Figure S19) showed that the Li||LFP cells with MCE consistently delivered higher discharge capacities than those with BE at all tested current densities, with the capacity difference becoming more pronounced at elevated rates.

FIGURE 5.

Electrochemical performance of lithium metal batteries with practical cathodes enabled by MCE. (a) Cycling performance of Li||LFP cells at 2 C, and corresponding (b) charge/discharge voltage profiles of Li||LFP cells at selected cycle numbers. (c) Rate capability of Li||NCM622 cells at different C‐rates and (d) relative discharge capacity enhancement of MCE compared with BE at different C‐rates, calculated as (Q _MCE – Q _BE) /Q _BE × 100%. (e) Cycling performance of Li||NCM622 cells at 4 C. (f) Cycling performance of Li||NCM622 cells at 10 C, and corresponding (g) the state of charge divided into CC/CV stages at selected cycle numbers.

To assess the capability of the electrolyte in high‐energy‐density systems, Li||NCM622 cells were assembled. The MCE‐based cells delivered superior rate performance over the entire tested range (0.5−5 C), consistently outperforming BE‐based cells (Figure 5c,d; Figure S20). At 5 C, where increased polarization tends to prematurely limit the accessible capacity in LMBs, MCE still delivered a discharge capacity 10.5% higher than that of BE, reflecting reduced polarization under high current conditions. To explore its practical applicability, long‐term cycling stability was evaluated at different rates. Across charging rates of 4, 8, and 10 C with the discharge rate fixed at 1 C (Figure 5e–g; Figures S21,S22), the Li||NCM622 cells with MCE exhibited higher discharge specific capacities and slower degradation. A clear divergence in long‐term cycling stability emerged at 4 C (Figure 5e), with the MCE‐based cells retaining 88.2% of their initial capacity after 250 cycles, compared to 72.3% for BE. Notably, even at 10 C, a condition typically associated with severe polarization and interfacial instability, the cells with MCE maintained excellent cycling stability. As shown in Figure 5f, after 500 cycles at 10 C, the MCE‐based cells can still deliver approximately 130 mAh g⁻¹, while BE cells dropped to 100 mAh g⁻¹. Using 80% capacity retention as the failure criterion, the MCE remained stable for 380 cycles, nearly doubling the lifetime of BE (190 cycles). Further insight was provided by the constant‐current/constant‐voltage (CC/CV) charging analysis (Figure 5g). The MCE‐based cells exhibited a higher proportion of CC charging across all cycle numbers, indicating that the applied charging current could be sustained for a longer duration under elevated current densities, and that charge transfer remained efficient and the electrode/electrolyte interface maintained low resistance. In contrast, BE cells showed a growing shift toward CV charging as cycling proceeded, reflecting increased polarization and worsening transport limitations. The cycling stability of the MCE‐based Li||NCM622 cells remains competitive among recently reported liquid electrolytes under comparable current densities (Table S3). To further evaluate practical feasibility, a Li||NCM622 pouch cell (∼10 mg cm⁻²) was tested (Figure S23), where the MCE‐based cell shows stable cycling behavior at 0.3 C. Overall, lithium metal batteries with practical cathodes show that the performance advantages enabled by MCE are consistent with the spatiotemporally continuous Li⁺ transport established above, thereby translating the stabilized lithium behavior into robust fast‐charging cycling.

3. Conclusions

In this work, a sulfonated MXene (Ti₃C₂(SO₃H)_x) colloid electrolyte was developed to address spatiotemporal discontinuity in Li⁺ transport under fast‐charging conditions, which limits the stability of lithium metal batteries. The ─SO₃H terminations alter the Li⁺ coordination environment, thereby lowering the desolvation barrier, while the conductive MXene sheets regulate the local potential and current distribution, collectively supporting a Li⁺ flux that is stably maintained during operation. This cooperative regulation mitigates the local amplification of transport driving forces under high current densities, leading to a more persistent and well‐regulated Li⁺ supply. As a result, Li||Li symmetric cells employing the MCE exhibit substantially improved interfacial stability, sustaining stable cycling for over 600 h at 1 mA cm⁻² and 1 mAh cm⁻² while remaining stable at high current densities. This enhanced stability is accompanied by compact and uniform lithium deposition on the anode surface after prolonged cycling. Accordingly, the MCE‐based cells exhibit outstanding rate capacity and cycling stability, retaining 90% capacity over 3000 cycles at 2 C in Li||LFP cells and maintaining over 80% capacity for more than 380 cycles at 10 C in Li||NCM622. These findings underscore the importance of enabling spatiotemporally continuous ion transport for achieving fast‐charging lithium metal batteries, and provide a clear transport‐oriented strategy for electrolyte design.

4. Experimental Section and Methods

The detailed experimental procedures are provided in the Supporting Information.

Funding

National Natural Science Foundation of China (No. 22378139), Guangdong Basic and Applied Basic Research Foundation (No. 2024A1515011370), Guangzhou Applied Basic Research Foundation (No. 2024A04J3037), and Ph.D Innovation Fund of GAC Group (No. XGA‐FD‐5).

Conflicts of Interest

The authors declare no conflicts of interest

Supporting information

Supporting File: smll73862‐sup‐0001‐SuppMat.docx.

Data Availability Statement

Data will be made available on request.