Anion‐Polarity‐Induced Interfacial Reconstruction Directs Crystallographic Texture for Highly Reversible Sn Anodes

ABSTRACT

Controllable Sn electrodeposition is crucial for achieving highly reversible Sn metal batteries, but remains a significant challenge. Here, preferential exposure of Sn(211) planes with a uniform and compact morphology is achieved via an anion‐polarity‐driven interfacial regulation strategy during plating and stripping. Strong and crystallographically selective adsorption of CF₃SO₃ ⁻ on Sn effectively modulates deposition behavior. Owing to its large intrinsic dipole, CF₃SO₃ ⁻ establishes a pronounced surface‐normal interfacial dipole field on Sn(211), which suppresses growth along this plane while accelerating deposition on competing orientations, thereby driving the gradual evolution toward dominant Sn(211) exposure. Consequently, the reconstructed Sn anode delivers outstanding performance, including 99.88% Coulombic efficiency over 10,200 cycles at 10 mA cm⁻² in Sn//Cu cells, ultralong cycling stability of 3456 h (∼4.8 months) in Sn//Sn symmetric cells, and 71% capacity retention after 20,000 cycles in full cells. The dipole‐based interfacial descriptor proposed in this work correlates anion polarity with crystallographic texture evolution in metal electrodeposition, providing a rational route toward controllable electrochemical reconstruction for high‐performance metal batteries.

Anion‐dipole‐driven interfacial regulation enables preferential Sn(211) orientation for highly reversible aqueous Sn batteries.

1. Introduction

Acidic aqueous tin (Sn) metal batteries have attracted considerable attention owing to their intrinsic safety, low cost, and favorable properties of metallic Sn [1, 2, 3, 4, 5, 6]. However, in acidic electrolytes, the high proton activity drives the hydrogen evolution reaction and continuous dissolution‐redeposition of Sn, resulting in a highly dynamic and unstable interfacial structure [6, 7]. Specifically, the electric double layer (EDL) and local interfacial microenvironment undergo persistent fluctuations, leading to highly disordered nucleation and growth across multiple crystallographic planes [8]. Moreover, the intrinsically small surface‐energy anisotropy of β‐Sn further weakens the thermodynamic driving force for preferential orientation [9, 10, 11]. Consequently, oversized and loosely packed polyhedral Sn particles are formed and unevenly distributed across the electrode. These particles exhibit poor adhesion to the current collector and readily detach during cycling to form electrochemically inactive “dead Sn”, thereby significantly lowering the Coulombic efficiency [1, 2]. Therefore, achieving uniform and dense deposition under acidic conditions remains a key scientific challenge for enabling highly reversible Sn batteries.

Whether a dense and homogeneous Sn deposition can be achieved is fundamentally determined by the evolution of the electrode/electrolyte interfacial structure during electrochemical deposition [12, 13, 14]. This interfacial structure is commonly described by the EDL, consisting of solvent molecules, cations, and anions [15, 16]. The organization of these interfacial elements will define the local electrostatic environment, thereby modulating desolvation barriers and nucleation kinetics, and further influencing atom incorporation and migration pathways, ultimately shaping the plating and stripping processes [17]. It is demonstrated that Sn deposition can be appropriately regulated by tuning solvation structures, cation coordination environments, and mass transport [1, 2, 3, 4, 5]. However, the specific role of anions in these processes has always been qualitatively discussed. Moreover, in strongly acidic electrolytes, rapid interfacial restructuring under highly reactive conditions further obscures anion‐specific pathways and makes their crystallographically differentiated interfacial behavior particularly difficult to resolve.

At the molecular scale, the local interfacial electric field is influenced not only by net charge distribution but also by the spatial orientation of interfacial species, particularly dipolar arrangements [18]. Anions with intrinsic dipole moments may adopt crystallographically distinct adsorption configurations and orientations, giving rise to facet‐dependent dipolar contributions that modulate the local reaction microenvironment and influence nucleation site selection and early growth pathways [19]. The distinct electrochemical reversibility and deposition morphology of Sn anodes have been observed with different anions [2, 3]. Generally, it is interpreted from qualitative perspectives within solvation‐ and transport‐related frameworks [3, 11], while the spatial expression of anions at the interface and the resulting crystallographically differentiated reaction conditions remain insufficiently resolved. Importantly, a quantitative descriptor capable of linking anion dipole characteristics to interfacial electric‐field reconstruction and facet‐dependent reactivity is still lacking, limiting deeper mechanistic understanding of morphology evolution in acidic Sn deposition.

In this research, an anion‐polarity‐driven interfacial regulation strategy is proposed to achieve compact and preferentially oriented Sn deposition in acidic aqueous electrolytes for highly reversible Sn metal batteries (Scheme 1). Moreover, a dipole‐based interfacial descriptor is further defined to capture the plane‐specific normal dipole contributions and adsorption densities of anions to quantitatively describe interfacial dipole organization biasing crystallographic texture evolution. Anion dipoles contribute to interfacial polarization, thereby reshaping the interfacial electric field and modulating the near‐interfacial Sn²⁺ transport and reaction flux. Crystallographically selective anion alignment and adsorption configurations, defined by their normal dipole components and interfacial orientations, establish orientation‐dependent local electric fields that guide the differential evolution of Sn crystal planes and ultimately promote the development of preferred Sn(211) texture. Consequently, the anion‐engineered Sn anodes deliver exceptional cycling reversibility in Sn//Sn symmetric cells (stable plating/stripping over 3456 h), Sn//Cu half cells (99.88% Coulombic efficiency over 10,200 cycles at 10 mA cm⁻²), and Sn//PANI full cells (71% capacity retention after 20,000 cycles). This work provides molecular‐scale insights into anion‐polarity‐directed interfacial engineering for precise facet control and stable metal deposition in acidic aqueous systems.

SCHEME 1.

Schematic illustration of anion‐polarity‐regulated interfacial dipole structures governing facet‐selective Sn deposition, where CF₃SO₃ ⁻ exhibits a strong intrinsic dipole, CH₃SO₃ ⁻ shows a moderate dipole, and SO₄ ²⁻ possesses a negligible dipole contribution. On different crystallographic planes, CF₃SO₃ ⁻ and CH₃SO₃ ⁻ display pronounced plane‐dependent normal dipole components (µ_z2 > µ_z1 > µ_z4 > µ_z3) and adsorption tilt angles (θ₂ > θ₁ > θ₄ > θ₃) with distinct adsorption densities, whereas SO₄ ²⁻ exhibits non‐selective adsorption and an almost negligible normal dipole component.

2. Result and Discussion

Electrodeposition morphology and crystallographic orientation are primarily governed by the interfacial electric‐field distribution, which in turn depends on the physicochemical characteristics of the electric double layer (EDL) [12, 14, 20]. While cationic additives have been extensively employed to modulate the EDL through combined electrostatic shielding and steric effects [21, 22], thereby enabling uniform and compact electrodeposition, the interfacial regulation by anionic species, especially in acidic systems, remains comparatively underexplored. In concentrated strong‐acid electrolytes, the EDL consists of a dense inner Helmholtz plane (IHP), predominantly occupied by hydronium ions (H₃O⁺) and anions, along with an outer diffuse layer [20, 23]. Anions possessing intrinsic dipole moments can perturb local interfacial electric fields, reorganize the interfacial hydration layer, and influence ion transport dynamics [8, 24, 25, 26]. Such dipole‐driven perturbations manifest as interfacial polarization, reshaping local electric fields and thereby influencing nucleation behavior and crystallographic orientation [27]. At the molecular scale, interfacial polarization originates from the collective alignment of anion dipoles, with the polarization magnitude governed by dipole surface density and orientation along the surface normal. According to electrostatic boundary conditions, interfacial polarization can be approximately described by a local dipole‐driven interfacial index (LDDI), defined as:

$$\mathbf{LDDI}=\mathit{n}\times 〈{\mathit{\mu }}_{\mathit{z}}〉$$

where n represents the effective anion adsorption density, and <µ _z> denotes the average surface‐normal dipole moment component, reflecting the statistical distribution of adsorption configurations and molecular orientations. The LDDI serves as a simplified phenomenological descriptor for evaluating the anion‐dipole‐mediated interfacial polarization and local electric‐field reconstruction. Variations in LDDI among different crystal planes reveal crystallographically selective anion adsorption, distinct interfacial polarization, and the resulting local interfacial electric‐field reconstruction. Such interfacial reconstructions modulate the free‐energy landscape of ion desolvation and nucleation, making the nucleation barrier sensitive to crystallographic orientation [28, 29, 30]. Anions characterized by high LDDI on specific crystal planes can form interfacial dipole configurations with pronounced surface‐normal electric‐field components, which can promote directional ion desolvation and local nucleation while hindering adatom diffusion and incorporation, thereby retarding growth along these orientations and preserving such planes as the dominant exposed facets. Importantly, by quantifying anion‐dipole‐mediated interfacial polarization, LDDI may serve as a comparable descriptor of interfacial electrostatic reconstruction, offering an electrostatic perspective on orientation‐dependent crystal growth behavior.

In this research, SnSO₄, Sn(CH₃SO₃)₂ (Sn(MSA)₂), and Sn(CF₃SO₃)₂ (Sn(OTf)₂), containing SO₄ ²⁻, MSA⁻, and OTf⁻ anions, respectively, are selected to systematically elucidate the intrinsic correlation among anion polarity, solvation structure, and local dipole organization in electrolytes, which collectively define the molecular‐scale genesis of the LDDI governing interfacial polarization. Anion polarity dictates the electronic structure and electrostatic landscape that shape its solvation behavior. As shown in Figure 1a, the electrostatic potential (ESP) maps reveal distinct charge localization patterns. SO₄ ²⁻ exhibits the strongest localized negative ESP with tetrahedrally concentrated charge [29, 31], whereas OTf⁻ displays the highly delocalized ESP arising from CF₃‐induced charge dispersion [20, 32]. By comparison, MSA⁻ shows an intermediate ESP distribution with sulfonate‐polarized yet methyl‐moderated distribution [20], indicating partial charge delocalization. These electrostatic features govern the microscopic electrostatic environment that dictates local dipole organization in the electrolyte. The optimized solvation configurations further reveal how polarity differences give rise to distinct coordination structures (Figures 1b and S1). In SO₄ ²⁻ electrolyte, the solvation shell is dominated by strong hydration with limited anion participation, resulting in a compact and rigid local environment. The MSA⁻ system features mixed coordination of water and sulfonate groups, leading to moderate structural flexibility. By contrast, the OTf⁻ electrolyte forms a relaxed and highly polarizable solvation shell owing to the weakly coordinating and charge‐delocalized nature. These solvation features indicate an anion‐polarity‐governed transition from rigid hydration to flexible polarization, which is further supported by statistical analyses of molecular dynamics (MD) trajectories. As shown in Figure 1c, the proportion of contact ion pairs (CIPs) follows an unconventional hierarchy of OTf⁻ > MSA⁻ > SO₄ ²⁻, which deviates from the coordination‐strength sequence typically expected from the chemical binding affinities [28, 33]. Consistently, radial distribution functions (RDFs) demonstrate that the spatial proximity between ions in solution is governed by anion polarity rather than by static coordination strength [9]. In SnSO₄ electrolyte, the Sn–O_water peak shows high intensity and dominant water coordination (Figure 1d), revealing a rigid and water‐rich hydration cage with minimal anion penetration. In comparison, Sn(MSA)₂ and Sn(OTf)₂ electrolytes exhibit a progressive decrease in Sn–O_water peak intensity along with an increase in Sn–O_anion peak intensity (Figure 1e,f), indicating competitive substitution of H₂O by MSA⁻ and OTf⁻ in the first solvation shell, as further resolved by the coordination numbers (CN). Anion coordination rises from about 1.15 in SO₄ ²⁻ to roughly 2.23 in MSA⁻ and 2.68 in OTf⁻, while the corresponding water coordination decreases from around 5.38 H₂O in SnSO₄ electrolyte to about 4.40 and 3.82 H₂O in MSA⁻ and OTf⁻ systems. The bulky and charge‐delocalized OTf⁻ weakens local electrostatic screening around Sn²⁺ and enables a flexible solvation shell for transient and reversible ionic proximity, whereas the highly symmetric and strongly hydrated SO₄ ²⁻ stabilizes a rigid water cage suppressing direct ionic contact. This weakened water coordination around Sn²⁺ is further supported by the storage behavior of the electrolytes, with the Sn(OTf)₂ electrolyte remaining clear and retaining nearly unchanged ionic conductivity after 30 days, whereas the Sn(MSA)₂ and SnSO₄ electrolytes become turbid and show changes in ionic conductivity upon prolonged storage (Figure S2). These results highlight the contribution of anion polarity to solvation‐shell relaxation.

FIGURE 1.

Anion polarity effects on solvation structure. (a) ESP of SO₄ ²⁻, MSA⁻, and OTf⁻. (b) MD snapshots of Sn²⁺ solvation structure in different electrolytes. (c) CIP fractions in different electrolytes. RDFs and coordination numbers of (d) SnSO₄, (e) Sn(MSA)₂, and (f) Sn(OTf)₂ electrolytes. (g) Residence time of H₂O in the first solvation shell. (h) ¹H NMR spectra in different electrolytes. (i) Anion dipole average tilt angles and dipole increments in different electrolytes. (j) The fitted Raman spectra of O−H stretching vibration in different electrolytes.

The flexibility of solvation shell is further supported by the hydrogen‐bond network features and the residence time (τ) of water molecules. A dense hydrogen‐bond network and a long τ of first‐shell water molecules are observed in SnSO₄ electrolyte (Figures S3 and 1g). In Sn(OTf)₂ electrolyte, the number of hydrogen bonds is reduced and τ is shortened, thereby enabling frequent exchange and reorganization within the shell and reflecting pronounced solvation dynamics and orientational freedom. ¹H NMR measurements further confirm the weakening of hydrogen‐bond interactions in Sn(OTf)₂ electrolyte, as evidenced by an upfield shift of proton resonances (Figure 1h), indicating a flexible solvation shell [33]. This structural flexibility enables orientational rearrangement of dipoles within the solvation environment, as reflected by the solvation‐scale dipole distributions. The average solvation tilt angle (θ_sol) and the increment in the normal dipole component (Δµ_{s ol}) reveal that Sn(OTf)₂ electrolyte develops an ordered dipole configuration with a pronounced preferential orientation (Figure 1i), whereas SnSO₄ electrolyte exhibits a dispersed dipole arrangement and Sn(MSA)₂ electrolyte displays intermediate behavior. Consistent with the solvation dynamics and dipole reorganization, FT‐IR and Raman analyses further substantiate the anion‐polarity‐dependent coordination environment and hydrogen‐bond configuration of the solvation shell (Figures S4 and 1j). Accordingly, the anion with pronounced molecular polarity and delocalized charge distribution imparts strong orientational tunability to the solvation structure and promotes interfacial normal dipole accumulation.

The crystallographic dependence of anion polarity effects was inferred from the interfacial structure on Sn(211) and Sn(101) planes, which differ in surface openness and coordination environments [7]. As shown in Figure 2a, MD snapshots of Sn(211) and Sn(101) interfaces display that increasing anion polarity and electron delocalization consistently reshape the interfacial enrichment layer, with SO₄ ²⁻ forming a compact region near the surface, MSA⁻ adopting a diffuse configuration, and OTf⁻ establishing an enrichment layer positioned father from the interface. The resulting structural ordering correlates with a reduction in differential capacitance (Figure S5), indicating decreased interfacial permittivity and enhanced dipole‐driven interfacial polarization as anion polarity increases [2, 4]. This polarity‐dependent interfacial reorganization is also reflected in the depth‐resolved anion density distributions (Figure S6), which show progressively strengthened anion accumulation with increasing polarity and exhibit crystal‐plane‐specific spatial characteristics on Sn(211) and Sn(101) planes, thereby establishing the structural framework for subsequent cation organization [24]. As anion polarity increases, the integrated cation density on both surfaces rises systematically, while the spatial position and extent of the cation‐enriched region differ between the two crystallographic planes (Figure 2b–d), thus shaping the distribution of the interfacial electric field [20]. Meanwhile, the adsorption tilt angles and normal dipole components reveal that polarity‐driven interfacial dipole alignment is strongly regulated by crystallographic plane (Figure 2e), as further supported by polarized ATR‐IR spectroscopy. A stronger surface‐normal S−O response is observed for OTf⁻ on Sn(211) than on Sn(101), whereas MSA⁻ shows a weaker contrast and SO₄ ²⁻ exhibits no well‐resolved positive dichroic response (Figure S7). Anions such as MSA⁻ and OTf⁻, possessing intrinsic dipole moments, preferentially adopt surface‐normal adsorption orientations, and Sn(211) plane consistently exhibits larger normal dipole components than Sn(101) plane, as its open atomic arrangement enhances intrinsic dipole projection onto the surface normal. By contrast, the nearly dipole‐free SO₄ ²⁻ displays only minimal dipole contribution on both planes, indicating the limited participation in interfacial dipolar alignment. With the largest intrinsic dipole moment, OTf⁻ exhibits the strongest normal dipole contribution on Sn(211), where polarity‐defined differences are most distinctly expressed. These observations indicate that polarity‐induced interfacial dipole structures manifest distinct crystal‐plane‐dependent normal alignment, consistent with crystallographically regulated LDDI behavior. The combined variations in normal dipole components and anion adsorption densities consistently yield higher LDDI values on Sn(211) than on Sn(101), with the amplification most pronounced for anions possessing stronger intrinsic dipole moments (Figure 2f and Table S1). Moreover, interfacial water density directly reflects the crystal‐plane‐dependent expression of interfacial polarization. As shown in Figure 2g, SnSO₄ electrolyte maintains the highest water density on both Sn(211) and Sn(101) planes, suggesting weak interfacial polarization, whereas Sn(MSA)₂ electrolyte exhibits moderate and comparable water densities on the two planes (Figure 2h). In contrast, Sn(OTf)₂ electrolyte displays a pronounced crystal‐plane asymmetry, with the water density on Sn(211) plane being significantly lower than that on Sn(101) plane (Figure 2i), revealing a more strongly polarized and more effectively desolvated interface on Sn(211) plane. Consistently, the effective penetration depth of hydrated protons decreases with increasing interfacial polarization (Figure 2j), demonstrating that anions with higher intrinsic polarity sustain a more concentrated and smoother interfacial electric‐field gradient. In addition, interfacial Raman spectra reveal attenuated O−H signals and sharpened OTf⁻ vibrational bands relative to the bulk electrolyte (Figure S8), further supporting the enhanced interfacial polarization and water depletion. This crystal‐plane‐dependent interfacial polarization is expected to govern the facet‐selective nucleation and growth behavior during Sn deposition.

FIGURE 2.

Anion‐regulated interfacial structure on distinct Sn(101) and Sn(211) planes. (a) Snapshots of interfacial electric double layer configurations in different electrolytes. Density profiles of Sn²⁺ on Sn(101) and Sn(211) planes in (b) SnSO₄, (c) Sn(MSA)₂, and (d) Sn(OTf)₂ electrolytes. (e) Average dipole tilt angles of anions on Sn(101) and Sn(211) planes in different electrolytes. (f) Anion adsorption densities, average surface‐normal dipole moment components, and LDDI values on Sn(101) and Sn(211) planes in different electrolytes. Density profiles of H₂O on Sn(101) and Sn(211) planes in (g) SnSO₄, (h) Sn(MSA)₂, and (i) Sn(OTf)₂ electrolytes. (j) Penetration depths of H₃O⁺ on Sn(101) and Sn(211) planes in different electrolytes.

To evaluate the influence of anion‐governed interfacial polarization on Sn nucleation and subsequent growth, Sn nucleation behavior was examined under different anion environments. As shown in Figure S9, Sn(OTf)₂ electrolyte exhibits the strongest nucleation response, with the lowest nucleation overpotential and the largest nucleation peak area. Chronoamperometry further reveals a brief initial current surge followed by rapid stabilization, indicating a dense population of simultaneously activated nuclei (Figure 3a) [11, 34]. By contrast, electrolytes containing low‐polarity anions show gradually rising currents, signifying dispersed nucleation and growth dominated by lateral diffusion [2]. According to the Scharifker‐Hills model [11, 35], all systems exhibit three‐dimensional instantaneous nucleation, with strongly polarized interfaces showing shorter times to the current maximum and higher peak currents, indicating higher instantaneous nucleation density, smaller nucleus size, and lower nucleation barrier (Figure 3b). Consistently, in situ EIS reveals the lowest charge transfer resistance (R _ct) and a smooth R _ct evolution in Sn(OTf)₂ electrolyte during the initial deposition stage, whereas the Sn(MSA)₂ and especially SnSO₄ electrolytes retain higher R _ct (Figure S10). These kinetic characteristics are clearly reflected in the subsequent morphological evolution. Under strong interfacial polarization, the deposition front remains smooth and continuous throughout plating (Figure 3c) [4, 36]. In contrast, weak interfacial polarization for low‐polarity anions leads to the early onset and rapid propagation of surface roughening and localized protuberances, driving heterogeneous growth, as further corroborated by the concurrent evolution of interfacial adsorption features. In situ Raman spectroscopy for high‐polarity anion systems shows progressive intensity evolution and subtle peak shifts of interfacial vibrational modes (Figures 3d and S11), indicating continuous reorganization of the interfacial environment [37 ^– 39]. By contrast, electrolytes containing low‐polarity anions display only marginal spectral evolution, consistent with weak interfacial polarization and minimal interfacial restructuring. These interfacial response differences are directly translated into distinct deposit morphologies. The high‐LDDI system yields a flat and uniform deposit composed of densely packed grains, whereas low‐LDDI systems generate coarse, porous, and structurally irregular deposits (Figures 3e and S12). Three‐dimensional surface reconstructions further visualize this contrast and reveal a narrow height distribution (R _a = 1.11 µm) for the high‐LDDI electrolyte, in contrast to the pronounced multiscale roughness and height fluctuations (R_a ≥ 2.15 µm) observed for the low‐LDDI electrolytes (Figure 3f).

FIGURE 3.

Sn electrochemical and deposition behaviors. (a) (I/I _m)²−(t/t _m) curves derived from chronoamperometry, and theoretical instantaneous and progressive nucleation profiles described by the Scharifker‐Hills model (2DP, 2DI, 3DI, and 3DP). (b) Density and radius of Sn nuclei at varying overpotentials. (c) In situ optical microscopy of Sn deposition. (d) In situ Raman spectra recorded during Sn plating and stripping processes. (e) SEM images of plated Sn at different capacities. (f) Three‐dimensional topographical reconstruction of Sn deposit. (g) XRD patterns of plated Sn at different capacities.

To elucidate the origin of the flat and compact morphology induced by high‐polarity anions, X‐ray diffraction (XRD) is performed on the corresponding tin deposits. As shown in Figure 3g, interfacial polarization plays a key role in directing the preferred crystallographic orientation. At a low plated capacity of 1 mAh cm⁻², the (211)/(101) intensity ratio (I₂₁₁/I₁₀₁) of Sn deposited from Sn(OTf)₂ electrolyte reaches 2.5, which is much higher than those of pristine Sn anode (0.8) and Sn(MSA)₂ (1.1) and SnSO₄ (1.0) electrolytes, indicating preferential Sn(211) deposition, consistent with the relative texture coefficient (RTC) analysis (Figure S13). The preferential exposure of Sn(211) plane is further confirmed by selected‐area electron diffraction (SAED). The Sn deposit derived from Sn(OTf)₂ electrolyte exhibits pronounced crystallographic alignment, as evidenced by discrete diffraction spots indexed to the (211) plane [7]. Conversely, deposits from Sn(MSA)₂ and SnSO₄ electrolytes exhibit diffraction patterns corresponding to multiple crystallographic orientations (Figure S14) [1]. As Sn deposition progresses in Sn(OTf)₂ electrolyte, the I₂₁₁/I₁₀₁ ratio rises to 8.9, which distinctly exceeds values below 6 observed in other electrolytes, indicating progressive reinforcement and stable expression of Sn(211) orientation. The orthogonal controls further show that this texture hierarchy cannot be consistently accounted for by acidity, salt concentration, ion transport, or corrosion behavior, but instead follows the anion‐dependent interfacial polarization (Table S2 and Figure S15). Strong interfacial polarization exerts directional control over Sn deposition by accelerating ion desolvation, synchronizing nucleation, and suppressing lateral‐diffusion‐induced roughening, preferentially stabilizing and exposing the Sn(211) plane and promoting the accumulation of a coherent crystallographic texture, ultimately yielding a dense, uniform, and structurally stable deposit (Figure S16). Insufficient interfacial polarization, however, provides little directional preference and leads to rough, randomly oriented structures. These findings demonstrate that local dipole‐driven interfacial polarization governs the crystallographically dependent LDDI response and macroscopic morphology of Sn deposition.

To confirm the positive effect of anion‐induced interfacial polarization on the reversibility of Sn plating/stripping behaviors, Sn//Cu asymmetric cells are examined under different conditions. At a low current density of 0.25 mA cm⁻², Sn//Cu asymmetric cell with Sn(OTf)₂ electrolyte delivers a stable plating/stripping process over 3245 cycles with an average Coulombic efficiency (Avg. CE) of 99.71%, outperforming Sn(MSA)₂ (1598 cycles, 99.06% Avg. CE) and SnSO₄ (476 cycles, 98.87% Avg. CE) electrolytes (Figure S17). When the current density increases to 1 mA cm⁻², the cell with Sn(OTf)₂ electrolyte sustains highly reversible cycling over 5866 cycles with an average CE of 99.85% at 0.5 mAh cm⁻² (Figure S18) and over 3280 cycles with an average CE of 99.85% at 1 mAh cm⁻² (Figure 4a). By contrast, the cells with Sn(MSA)₂ and SnSO₄ electrolytes quickly fail only after dozens of cycles with a low CE, primarily arising from coarse Sn particles and “dead Sn” formed during cycling. Meanwhile, the cell with Sn(OTf)₂ electrolyte exhibits stable plating/stripping overpotential without short‐circuiting during prolonged cycling, whereas cells using other electrolytes fail within only dozens of cycles (Figure S19). Importantly, Sn(OTf)₂ electrolyte displays a lower initial nucleation overpotential (∼3.3 mV) than that of other electrolytes (∼4.2 mV), suggesting a reduced barrier to initial nucleation (Figure S20) [40]. As expected, the Sn//Cu cell in Sn(OTf)₂ electrolyte demonstrates excellent high‐rate Sn plating/stripping reversibility, as evidenced by a consistently high CE over the range of 1 to 10 mA cm⁻². By comparison, the cell in Sn(MSA)₂ electrolyte shows noticeable CE fluctuation, while that in SnSO₄ electrolyte rapidly fails with severe CE fluctuation (Figures 4b and S21). At 5 mA cm⁻², the cell in Sn(OTf)₂ electrolyte delivers exceptional stability with an average CE of 99.96% over 15,700 cycles (Figure 4c), surpassing other electrolytes by >100‐fold in cycle life and by >0.7% in CE. Amazingly, an average CE of 99.88% can be achieved even over 10,200 cycles at 10 mA cm⁻² (Figure 4d), which is superior to previously reported acidic aqueous Sn batteries (Figure 4e and Table S3) [1, 2, 5, 9, 11], further demonstrating the effectiveness of anion‐induced high‐LDDI interface in improving reversibility of Sn anodes. As indicated by the side‐reaction results, the suppression of hydrogen evolution reaction (HER), Sn dissolution/corrosion, or self‐discharge does not necessarily correspond to excellent electrochemical performance (Figure S22). This highlights the positive role of the anion‐induced interfacial environment in promoting highly reversible Sn plating/stripping, with variations in side reactions emerging as coupled consequences.

FIGURE 4.

Electrochemical performance of Sn anodes in different electrolytes. (a) Coulombic efficiency of Sn//Cu cells at 1 mA cm⁻² and 1 mAh cm⁻². (b) Coulombic efficiency of Sn//Cu cells at 1 mAh cm^‒2 under different current densities. Coulombic efficiency of Sn//Cu cells at (c) 5 mA cm⁻² and (d) 10 mA cm⁻² with 1 mAh cm⁻², with the Sn(OTf)₂ results taken from representative Sn//Cu cells rather than the best‐performing cells. (e) Comparison of average CE of Sn//Cu cells with recently reported works [1, 2, 5, 9, 11]. (f) Cycling performance of Sn//Sn symmetric cells at 1 mA cm⁻² and 1 mAh cm⁻², with the Sn(OTf)₂ result taken from a representative Sn//Sn symmetric cell rather than the best‐performing cell. Cycling performance of Sn//Sn symmetric cells at 1 mA cm⁻² with a limited Sn supply at (g) DOD_Sn = 30% and (h) DOD_Sn = 60%.

Additionally, Sn//Sn symmetric cells are also assembled to evaluate the effect of anions on the cycling stability of Sn anode. As shown in Figure 4f, at 1 mA cm⁻² and 1 mAh cm⁻², the short‐circuit failure after cycling for 60 and 46 h can be observed in Sn(MSA)₂ and SnSO₄ electrolytes, respectively, whereas the cell with Sn(OTf)₂ electrolyte displays excellent cycling stability with a constant overpotential (∼10 mV) over 3456 h (∼4.8 months). With increasing current density and areal capacity, the structural reconstruction of Sn anode becomes pronounced, thereby weakening the anion regulation effect during prolonged cycling. Accordingly, the performance of symmetric cells is further investigated at 5 and 10 mA cm⁻². As shown in Figures S23 and S24, symmetric cells using Sn(MSA)₂ and SnSO₄ electrolytes suffer short‐circuit failure within ∼15 h at 5 mA cm⁻² and only a few hours at 10 mA cm⁻², whereas the cell with Sn(OTf)₂ electrolyte maintains stable voltage hysteresis over 472 h at 5 mA cm⁻² and over 161 h even at 10 mA cm⁻². Furthermore, the rate capability is evaluated at current densities from 0.5 to 10 mA cm⁻² with an areal capacity of 1 mAh cm⁻² (Figure S25), corresponding to a 6% depth of discharge (DOD). The anion‐induced high‐LDDI interface significantly improves the rate capability. The sudden voltage drops arising from short‐circuiting induced by coarse Sn particles are observed at 5 mA cm⁻² in other electrolytes. By contrast, a stable overpotential without noticeable fluctuation is maintained even at 10 mA cm⁻² in Sn(OTf)₂ electrolyte. Furthermore, under limited Sn supply at 1 mA cm⁻² with a DOD of 30% (Figure 4g), the symmetric cell in Sn(OTf)₂ electrolyte sustains stable cycling for 242 h, whereas cells in other electrolytes fail within only ∼20 h. Even at a high DOD of 60%, the symmetric cell in Sn(OTf)₂ electrolyte still maintains stable cycling over 310 h (Figure 4h), far longer than that in other electrolytes.

To further clarify the role of anions in regulating the Sn plating/stripping behaviors during cycling, the surface characteristics of cycled Sn anodes were investigated. As shown in Figure S26, cycled Sn anodes in different electrolytes exhibit distinct (211) peak intensities, indicating the superior stability of anion‐induced high‐LDDI interface in regulating crystallographic orientation. Compared with other anions, OTf⁻ markedly enhances the (211) diffraction peak intensity of cycled Sn anodes. After 100 cycles, the (211) texture emerges as the dominant crystallographic orientation in Sn(OTf)₂ electrolyte. The corresponding RTC analysis and pole figures further confirm the OTf⁻‐regulated preferential orientation (Figures S27 and S28). Additionally, the morphology evolution of Sn anodes during cycling is also investigated. As shown in Figure S29, cycled Sn anodes in Sn(MSA)₂ and SnSO₄ electrolytes develop visible protuberances and scalloped pitting cavities after only 5 cycles, followed by the evolution of loosely packed granular porous structures after 50 and 100 cycles due to uneven deposition, together with pronounced surface greying and optical inhomogeneity after cycling (Figure S30). By contrast, cycled Sn anode in Sn(OTf)₂ electrolyte retains a flat and compact morphology even after 100 cycles, while maintaining a more uniform optical appearance and metallic luster, indicating the excellent resistance to structural reconstruction enabled by the high‐polarity anion during prolonged cycling. More importantly, surface chemical analyses of cycled Sn electrodes reveal no evident S‐ or F‐related signals (Figures S31 and S32), indicating that the observed anion‐dependent surface evolution is not dominated by an anion‐derived passivation layer, but is more consistent with differences in the anion‐regulated interfacial environment during repeated Sn plating/stripping.

The interfacial polarization induced by anion dipoles, reflected by the LDDI as a conceptual strength indicator, exerts a positive influence on the deposition morphology and cycling stability of Sn anodes. To further delineate this influence, interfacial kinetic features and phase‐field characteristics are addressed. The Nyquist plots reveal a reduced charge‐transfer impedance with a contracted mid‐frequency arc for the high‐LDDI interface (Figure S33), indicating accelerated ion desolvation and charge transfer with lowered interfacial energy barriers [36]. Moreover, the corresponding charge‐transfer resistance remains low even after prolonged cycling (Figure S34), supporting the sustained kinetic activity and reversibility of the high‐LDDI interface. By contrast, the low‐LDDI interface maintains substantially high resistance, suggesting persistent kinetic limitations and hindered interfacial ion transport. Furthermore, the distribution of relaxation times (DRT) analysis provides a high‐resolution and pathway‐specific insight into the interfacial kinetic behavior (Figure 5a) [41]. At the high‐LDDI interface, the mid‐ and sub‐mid‐frequency relaxation modes associated with ion desolvation and early nucleation shift toward shorter characteristic times, indicating a compressed temporal window and a more temporally concentrated nucleation process [42]. By contrast, the low‐LDDI interface exhibits relaxation modes with broader profiles and elevated intensities, suggesting pronounced kinetic restrictions arising from the coupled diffusion‐reaction constraints at the interface. This kinetic divergence further demonstrates that anion‐mediated interfacial reconstruction modulates not only the overall rate of interfacial charge transfer and desolvation but also the temporal organization of nucleation and growth, offering a kinetic rationale for the observed crystallographic texture evolution.

FIGURE 5.

Anion‐dependent crystal‐plane‐selective interfacial kinetics. (a) DRT analysis in different electrolytes with τ₁, τ₂, τ₃, and τ₄ corresponding to fast desolvation and rearrangement, critical nucleus formation, early 3D growth, and diffusion‐layer mass transport. Finite‐element‐simulated (b, c) interfacial potential distributions and (d, e, and f) surface Sn²⁺ concentration evolution on different crystallographic planes.

In addition, finite‐element simulations were performed to further examine the spatially resolved kinetic differences among crystallographic planes [26, 43, 44]. The high‐LDDI interface does not generate uniform field intensities (Figure 5b,c) or ion‐flux distributions (Figures S35–S38) across different planes but instead displays pronounced crystallographic dependence. Regions adjacent to Sn(101) plane exhibit higher normal potential gradients and more focused cation fluxes, which collectively promote rapid nucleation and three‐dimensional growth. In contrast, regions near Sn(211) plane present weaker normal gradients and reduced flux convergence, resulting in a slower and more uniform advancement of the growth front. At the low‐LDDI interface, crystallographic selectivity is strongly attenuated because the weakened normal field fails to impose effective plane differentiation, resulting in stochastic growth trajectories and reduced interfacial stability. The intensified localization of field and flux at Sn(101) under high‐LDDI conditions directs a larger fraction of tin ions to react and accumulate on this plane (Figure 5d–f), accelerating its advance while the more slowly growing Sn(211) plane remains preferentially exposed. These results demonstrate that a high LDDI not only accelerates interfacial reaction kinetics overall but also establishes a pronounced plane‐dependent kinetic asymmetry, enabling the Sn(101) surface to accumulate material more rapidly and extensively than the Sn(211) surface, whereas such differentiated kinetic driving is absent under low‐LDDI conditions. Overall, these findings demonstrate that the anion‐induced interfacial dipole field establishes a pronounced kinetic asymmetry between crystallographic orientations, driving rapid adatom accumulation and overgrowth on Sn(101) so that this fast‐advancing orientation becomes buried during subsequent deposition, whereas the slower and more stable Sn(211) plane remains exposed and ultimately dominates the final texture.

Finally, the effectiveness of anion‐induced high‐LDDI interface is further demonstrated by the Sn//polyaniline (Sn//PANI) full cells. The Sn//PANI full cells display similar redox features (Figure 6a), whereas the cell with Sn(OTf)₂ electrolyte exhibits lower electrochemical polarization than those employing other electrolytes, arising from the enhanced reversibility of Sn²⁺/Sn conversion enabled by anion‐mediated interfacial optimization [45]. At 0.5 A g⁻¹, the full cell with Sn(OTf)₂ electrolyte delivers a capacity retention of 99.6% over 100 cycles (Figure S39), outperforming the cells with other electrolytes. Meanwhile, the full cell with Sn(OTf)₂ electrolyte exhibits the weakest self‐discharge behavior, retaining 87.9% of its initial capacity after resting for 24 h in the charged state (Figure 6b–d), which is markedly higher than that of Sn(MSA)₂ (77.7%) and SnSO₄ (75.1%), demonstrating efficient suppression of spontaneous Sn dissolution, as further supported by XRD analysis. Compared with other electrolytes, cycled Sn anode in Sn(OTf)₂ electrolyte preserves a preferred (211) orientation (Figure 6e), further confirming the regulatory effect of high‐polarity anion on Sn deposition. In contrast, abundant loose and isolated particle structures are observed on cycled Sn anodes in Sn(MSA)₂ and SnSO₄ electrolytes (Figure 6f). As a direct consequence of the preferential Sn deposition, Sn anode in Sn(OTf)₂ electrolyte remains dense and compact even after 100 cycles.

FIGURE 6.

Electrochemical performance of Sn//PANI full cells. (a) CV curves of full cells at a scan rate of 1 mV s⁻¹. Self‐discharge behaviors of full cells in (b) SnSO₄, (c) Sn(MSA)₂, and (d) Sn(OTf)₂ electrolytes. (e) XRD patterns and (f) SEM images of cycled Sn anodes from full cells. Cycling performance of full cells at (g) 1 A g⁻¹ and (h) 5 A g⁻¹, with the Sn(OTf)₂ results taken from representative full cells rather than the best‐performing cells. (i) Cycling performance of pouch cell in Sn(OTf)₂ electrolyte at 0.5 mA cm⁻².

The cycling performance is further evaluated under practical conditions by shortening the charging time. The full cells with Sn(MSA)₂ and SnSO₄ electrolytes display pronounced capacity decay after several hundred cycles, whereas the cell with Sn(OTf)₂ electrolyte sustains a high reversible capacity with an excellent capacity retention of 71% after 20,000 cycles at 1 A g⁻¹ (Figure 6g). Notably, even at a high current density of 5 A g⁻¹, a capacity retention of 51% is still maintained after 10,000 cycles (Figure 6h). Furthermore, the practical feasibility of the anion‐regulation strategy is demonstrated in pouch cells operated under lean‐electrolyte conditions. As shown in Figure S40, the pouch cell with Sn(OTf)₂ electrolyte displays charge‐discharge profiles comparable to those of the coin cell, indicating consistent electrochemical behavior upon scale‐up. The pouch cell delivers a high reversible capacity of 138 mAh g⁻¹ with an average CE of 99.9% over 700 cycles at 0.5 mA cm⁻² (Figure 6i), whereas pouch cells with Sn(MSA)₂ and SnSO₄ electrolytes fail rapidly within only a few cycles (Figure S41). These results further confirm the superior performance of Sn metal batteries enabled by the anion‐induced high‐LDDI interface. Moreover, this anion‐regulation strategy effectively couples solution‐phase thermodynamics with solid‐phase crystallization kinetics, thereby establishing a generalizable principle applicable to other p‐block metals in harsh electrolytes without complex additives or specialized substrates.

3. Conclusion

In summary, an anion‐polarity‐regulated interfacial dipole mechanism is established that maps a coherent multiscale pathway from molecular polarity, through interfacial reconstruction, local electric‐field reconfiguration, and crystal‐plane‐dependent Sn²⁺ reaction‐flux differentiation, to the preferred crystallographic texture during Sn electrodeposition. The oriented anion‐dipole response reorganizes the interfacial environment and local electric‐field distribution, giving rise to crystal‐plane‐dependent Sn²⁺ reaction‐flux differentiation that promotes rapid adatom accumulation and overgrowth associated with Sn(101) plane, whereas the kinetically slower and more stable Sn(211)‐oriented component persists during continued deposition and ultimately dominates the texture. This dipole‐mediated selective‐growth pathway not only directs crystallographic texture evolution but also promotes uniform deposition morphology, enhanced interfacial stability, and markedly improved electrochemical performance. This work reveals the important role of anion dipole regulation in governing the cross‐scale coupling between molecular interfacial chemistry and macroscopic crystallographic texture evolution, providing new mechanistic insights into metal electrodeposition.

Conflicts of Interest

The authors declare no conflict of interest.

Supporting information

Supporting File: adma74286‐sup‐0001‐SuppMat.docx.

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.