Controlled synthesis and interfacial properties of polyvinylidene fluoride based metal-fluoride surface treatments for high-nickel NCM cathodes

Abstract

Despite the advantage of reduced cobalt content, widely proposed Ni-rich cathode candidates face significant challenges related to thermal instability and structural deterioration under high-voltage conditions. In particular, when the Ni content exceeds 80% of the total transition metal in LiNi_xCo_yMn_1−x−yO₂, the H2–H3 phase transition accelerates material degradation through irreversible structural evolution during electrochemical cycling. A common irreversible reaction in Ni-rich cathodes arises from the weakening of Ni–O covalent bonds during Ni oxidation, which triggers the oxygen reduction reaction and subsequent gas evolution, thereby further accelerating material degradation under high-voltage operation. As such degradation predominantly occurs at the particle surface, numerous studies have employed surface-coating strategies based on metal cations that form strong bonds with oxygen in the host lattice and suppress oxygen release. However, coatings involving high-valence metal ions often reduce the initial capacity of the cathode material and hinder Li⁺ diffusion kinetics across the interphases. In this study, we present a straightforward surface-treatment strategy using fluorine anions derived from polyvinylidene fluoride (PVDF) to prevent gas evolution while maintaining both the capacity and lithium-ion diffusivity of a Ni-rich LiNi_0.96Co_0.035Mn_0.005O₂ cathode.

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Young-Sang Yu

Young-Sang Yu is an Associate Professor in the Department of Physics at Chungbuk National University. He earned his PhD in Materials Science and Engineering from Seoul National University, where he studied magnetic vortex dynamics in ferromagnetic nanostructures. Before joining Chungbuk National University, he worked at the Lawrence Berkeley National Laboratory and Advanced Light Source as a postdoctoral fellow and beamline scientist. His research combines synchrotron-based X-ray microscopy, tomography, and spectroscopy to investigate nanoscale phenomena in energy materials, catalysts, and magnetic systems.

1. Introduction

High energy density and efficient energy conversion have established lithium-ion batteries (LIBs) as the leading technology for modern energy storage systems. Among various cathode materials, layered NCM compounds (LiNi_xCo_yMn_1−x−yO₂) have emerged as predominant candidates due to their favorable balance between cost and electrochemical performance.^1,2 Due to the distinct electrochemical roles of each transition metal (TM) ion, Ni cations contribute predominantly to Li-ion storage capacity through reversible valence changes between +2 and +4, which underpins the enhanced capacity of high-Ni NCM cathodes.^3,4 In contrast, Co and Mn ions primarily improve electronic conductivity for better rate performance and crystallographic robustness, respectively. However, when the Ni content exceeds 80% of the total TMs, pronounced surface rock-salt phase formation and the development of microcracks driven by large volumetric changes during Li-ion (de)intercalation significantly deteriorate cycling stability and compromise safety performance.⁵ As a result, high-Ni NCM materials typically exhibit poor capacity retention, a drawback that becomes more severe under high-voltage and high-temperature conditions relevant to high-power-density applications.

The performance decline in high-Ni NCM cathodes is primarily driven by multiple degradations mechanism, including chemical dissolution,^6,7 phase transformations from the layered to rock-salt phase,^4,8 and mechanical damage,^9,10 all of which are closely associated with oxygen release from the host crystal lattice, leading to a weakened oxygen framework and an altered local chemical environment.^5,11 For example, deeper delithiation induced by higher Ni content creates additional vacancies in the Li-ion layer, leading to increased cation intermixing, structural instability, and TM dissolution.¹² Furthermore, oxygen released from cathode materials, often in the form of highly reactive singlet (¹O₂), can trigger electrolyte oxidation and decomposition reactions,^13,14 potentially initiating “rollover” effects that culminate in thermal runaway and, in extreme cases, battery cell failure or explosion.^15,16 Among the various degradation pathways induced by oxygen loss, phase-transformation-associated structural degradation is directly responsible for the diminished electrochemical performance of high-Ni cathodes.¹⁷ Increasing Ni content typically shifts the original phase transition behavior during (dis)charging toward highly anisotropic structural changes associated with the H2–H3 transition, thereby accelerating a degradation pathway from the layered (R3̄m, LiMO₂) phase to the spinel (Fd3̄m, LiM₂O₄) phase and ultimately to the electrochemically inactive rock-salt (Fm3̄m, MO) phase, particularly in surface-localized regions of the cathode particles.¹⁸ The formation of a NiO-type rock-salt phase not only signifies irreversible collapse of the layered structure required for Li (de)intercalation but also substantially increases the kinetic barrier for Li-ion diffusion, thereby reducing the attainable energy density of the overall cathode active material.^19–21

To achieve both high capacity and stable cycling performance in Ni-rich cathode materials, various strategies have been proposed, including surface coating,^22–25 gradient chemical composition,^26,27 cation and anion doping,^28,29 and morphology modifications.³⁰ Since oxygen loss and the associated structural degradation are primarily localized at the particle surface, surface stabilization through an appropriate coating layer can significantly decelerate degradation kinetics, thereby mitigating electrochemical degradation.^5,31 Among the commonly reported coating configurations,³² electrochemically inactive metal-fluoride coatings, such as AlF₃,^33,34 FeF₃,³⁵ and CaF₂,³⁶ are typically applied in either core–shell or ultra-thin coating geometries. However, thick coating layers, as in core–shell architectures, can reduce the initial capacity by limiting the fraction of active material and do not further enhance capacity retention.³⁴ Therefore, an ideal coating layer should uniformly cover the entire particle surface while maintaining minimal thickness. Achieving such ultrathin and conformal coatings, however, often requires highly precise synthesis control and carries the risk of introducing unexpected secondary phases.³⁷ Moreover, the thermal conversion of high-Ni precursors into layered cathode materials must be conducted under high oxygen partial pressure to oxidize Ni²⁺ to Ni³⁺.³⁸ Under such oxygen-rich conditions, controlling fluorine incorporation into lattice oxygen sites becomes highly challenging, which can deteriorate cathode performance through parasitic reactions among TMs, lithium, and fluorine species.^39,40

Another interfacial-stabilization strategy is to improve surface hydrophobicity through thin polymer coatings containing hydrophobic functional groups.^41–43 Such solution-processable coatings can form a continuous, conformal protective layer on Ni-rich cathode particles via relatively simple synthesis routes.⁴¹ By limiting exposure to ambient H₂O and CO₂, which readily react with Ni-rich layered oxides, the hydrophobic coatings suppress the reduction of trivalent Ni, the formation of rock-salt surface layers, and the accumulation of LiOH/Li₂CO₃ residues, thereby mitigating parasitic interfacial reactions.^41–43 However, polymer-based hydrophobic coatings may suffer from limited interfacial robustness and, when excessively thick, can impede Li-ion transport and reduce electrochemical performance.⁴³

In this study, we propose a synergistic strategy that combines the processing simplicity of solution-based polymer coating with the electrochemical advantages of metal-fluoride surface treatment. Specifically, a solution-deposited polyvinylidene fluoride (PVDF) layer is converted through a simple, cost-effective, and highly controllable solid-phase reaction into an ultrathin metal-fluoride coating on preformed high-Ni NCM particles. Our approach enables precise control over fluorine incorporation while avoiding interference with Li-ion diffusion into the layered precursor structure. The electrochemical performance of high-Ni NCM cathodes with optimized metal-fluoride coatings was confirmed through high-temperature and high-voltage cycling, as well as galvanostatic intermittent titration technique (GITT) electrochemical analysis. The underlying origins of the improved stability were systematically analyzed by investigating the formation and chemical nature of the surface fluorine layer using X-ray photoemission spectroscopy (XPS), X-ray absorption spectroscopy (XAS), and scanning transmission X-ray microscopy (STXM). This work will provide deeper insight into the qualitative role and spatial distribution of fluorine within cathode particles, offering a viable strategy to mitigate oxygen-release-related degradation in layered cathode materials.

2. Experimental section

2.1 Preparation of the metal-fluoride coated high-Ni NCM active material

High-Ni LiNi_0.960Co_0.035Mn_0.005O₂ (NCM) was synthesized from a co-precipitated hydroxide precursor, Ni_0.960Co_0.035Mn_0.005(OH)₂, prepared following a previously reported procedure.⁴⁴ The precursor was mixed with LiOH·H₂O in a molar ratio of 1 : 1.02, preheated at 400 °C for 2 h, and then calcined at 710 °C for 12 h under an oxygen atmosphere. For fluorine surface treatment, the as-synthesized NCM powder was mixed with PVDF in acetone, dried at 60 °C, pelletized at approximately 27 MPa for 5 min, and subsequently heat-treated at 350 °C under Ar. Under optimized conditions, PVDF served as a fluorine precursor to generate an ultrathin metal-fluoride-rich surface layer on the NCM particles. Unless otherwise noted, the PVDF content was fixed at 1 wt%. A schematic illustration of the treatment process is shown in Fig. S1. Detailed synthesis procedures are provided in the SI.

2.2 Electrochemical tests

2032-type Li/NCM coin half-cells were assembled using lithium metal as the counter electrode and high-Ni NCM as the working electrode, with 1 M LiPF₆ dissolved in EC/DMC/EMC (1 : 1 : 1 by volume) as the electrolyte. Cathodes were prepared from uncoated or fluorine-coated (F-coated) high-Ni NCM powders using acetylene black and PVDF binder. The electrochemical performance of both samples was evaluated at various temperatures and cut-off voltages. To assess the effect of the fluorine surface treatment under harsh conditions, cycling tests were performed at 40 °C over a voltage window of 2.7–4.5 V. Electrochemical impedance spectroscopy (EIS) and galvanostatic intermittent titration technique (GITT) measurements were additionally conducted to monitor the evolution of Li⁺ diffusion kinetics during cycling. Detailed electrode fabrication, cell assembly, EIS, and GITT conditions are provided in the SI.

2.3 Material characterization

The particle surface morphology and elemental distribution of the synthesized cathodes were examined using field-emission scanning electron microscopy (FE-SEM, Tescan Clara, Tescan) coupled with energy-dispersive spectroscopy (EDS). ImageJ software⁴⁵ was used to quantitatively analyze morphological features at both the secondary- and primary-particle levels. The crystallographic information of the calcined products was investigated by powder X-ray diffraction (XRD, Bruker, D8 Advance) using Cu Kα radiation (λ = 1.54178 Å) operated at 40 kV and 40 mA. The diffractograms were collected in the 2θ range of 10–70° with a step width of 0.02° and a scan rate of 2° min⁻¹. Chemical analysis was further complemented using XPS (Thermo Fisher Scientific) with Al Kα radiation. For depth-dependent chemical analysis, depth-profiled XPS measurements were additionally performed using a PHI 5000 VersaProbe (ULVAC-PHI) equipped with Ar-ion sputtering.

2.4 Chemical state mapping via X-ray microscopy

Scanning transmission X-ray microscopy (STXM) was used to investigate the fluorine chemistry of F-coated NCM. F K-edge image stacks were acquired at Beamline 11.0.2.2 of the Advanced Light Source (ALS), Lawrence Berkeley National Laboratory,^46,47 using fragmented primary particles prepared from secondary particles. The incident photon energy near the F K-edge was calibrated using reference gases and a metallic thin-film standard.⁴⁸ Transmitted X-ray images were converted to optical-density (OD) maps and aligned with sub-pixel accuracy using an iterative registration procedure.⁴⁹ The fluorine chemical states were identified by comparison with reference spectra of LiF and PVDF.^50,51 To further examine the radial distribution of F-containing species together with the distance-dependent spectral variations of O species and Ni oxidation states, additional STXM measurements were performed at the 10A1 beamline of Pohang Light Source II (PLS-II)⁵² using focused ion beam (FIB)-milled lamellae prepared from pristine F-coated NCM secondary particles. Detailed acquisition parameters and data-processing procedures are provided in the SI.

3. Results and discussion

The surface morphologies of both uncoated and F-coated NCMs are shown in the SEM images in Fig. 1a and b, respectively. Both secondary particles exhibit a spherical shape with a mean diameter of approximately 20 µm and are composed of gravel-shaped primary particles. The samples display similar primary particle size distribution and consistency (D_50,uncoated = 21.094 µm and span_uncoated = 0.663; D_50,F-coated = 20.159 µm and span_F-coated = 0.739, Fig. S2), indicating that the post-treatment to form the metal-fluoride coating does not affect the original morphologies of the synthesized high-Ni NCM. The primary morphological difference could be observed, possibly due to the surface coating, which is featured by a slightly roughened texture on the surface. The roughened surface can originate from either the formation of a fluorine-containing layer with nonuniform thickness or residual carbon left after PVDF decomposes in an inert atmosphere. The coating provides overall complete coverage, effectively isolating the particles from exposure to the electrolyte.

Fig. 1 SEM images of (a) uncoated and (b) F-coated NCM particles. (c) XRD patterns of uncoated (soft coral pink solid line) and F-coated (light blue solid line) NCM particles with peak positions of ideal stoichiometric LiNiO₂ (JCPDS card no. 09-0063, black lines). (d) Zoomed-in regions near the (003), (104), (012), and (110) diffraction peaks. Ni 2p XPS spectra for (e) uncoated and (f) F-coated samples. In Ni 2p XPS spectra, Ni²⁺, Ni³⁺, and satellite are deconvoluted and fitted.

Consistent with the morphology observations, the crystallographic characteristics of both and F-coated NCMs show analogous features in their XRD patterns (Fig. 1c). In the high-Ni NCM crystal, the intensity ratio between the (003) and (104) XRD peaks serves as a reliable indicator for the crystal quality because the ratio is closely related to cation mixing due to the similar ionic radii of Ni²⁺ and Li⁺.^53–56 Overall, I₍₀₀₃₎/I₍₁₀₄₎ values are higher than 1.9 for both samples, which indicates a well-formed layered structure with a very small degree of Ni²⁺/Li⁺ cation mixing (Fig. 1d).^57,58 In addition, the distinct separation of the (006)/(102) and (108)/(110) XRD peaks observed in both samples confirms a high degree of crystallinity.^58,59 In addition, it is worth noting that additional impurity peaks, including LiF near the (102) and (104) peaks, were not observed in F-coated NCM.⁶⁰ The XRD analysis results confirm that the bulk properties of the NCM remain unchanged even after the surface fluorine coating treatment. Consequently, the different electrochemical behaviors, which will be discussed later, can be attributed to variations in the surface characteristics.

As an alternative analysis method to examine the chemical characteristics of both samples, XPS was employed owing to its surface sensitivity,⁶¹ as shown in Fig. 1e and f. The XPS survey spectra of each sample offer an overview of the surface elemental composition changes induced by the fluorine coating treatment. While the nickel element ideally exists in a form close to Ni³⁺ in the high-Ni NCM, an unintended rock-salt NiO phase is often introduced, particularly near the secondary particle surface, due to nickel reduction despite the thermal treatment process under high oxygen partial pressure.³⁸ In Ni 2p spectra (Fig. 1e), the Ni 2p_3/2 and Ni 2p_1/2 peaks are observed at ∼855 eV and ∼873 eV, respectively.⁵³ In addition, the broad features at ∼860 eV and ∼880 eV are satellites of Ni 2p_3/2 and Ni 2p_1/2, respectively. These peaks provide insight into the surface nature of the NCM samples. The presence of the rock-salt phase generated during the synthesis process can be identified by two peaks corresponding to Ni²⁺ and Ni³⁺ in the Ni 2p XPS spectra.⁵³ In the F-coated NCM (Fig. 1f), although the Ni²⁺ (NiO) and Ni³⁺ peaks were observed to be similar to those of the uncoated NCM (Fig. 1e), the Ni 2p_3/2 spectrum became slightly broader. This broadening is ascribed to the emergence of a new peak near 858 eV, which likely originates from the formation of the NiF₂ phase.

To provide a more in-depth analysis of the spatial distribution of the fluorine species, depth-profiled XPS was conducted, as shown in Fig. S4. The estimated atomic concentration of fluorine decreases gradually with increasing etching time, indicating that the majority of fluorine species are confined to the outermost surface region and do not significantly penetrate into the NCM bulk crystal lattice.

To further reveal the distribution of F-containing species and the spectral evolution of O and Ni as a function of distance from the secondary-particle surface, the focused-ion-beam (FIB)-milled, F-coated NCM particles were investigated using STXM across the F K-, O K-, and Ni L-edges, as shown in Fig. 2 and S5. Details of the STXM measurements and analysis are discussed in the Experimental section and SI. The overall Ni²⁺ concentration throughout the entire particle region, obtained by linear combination fitting of Ni²⁺ and Ni³⁺ reference spectra, reveals a pronounced Ni²⁺ fraction of 35.83% in pristine F-coated NCM and serves as an indicator of stoichiometric deviation from the ideal layered framework, as shown in Fig. S5a.

Fig. 2 (a) Optical density (OD) image of F-coated NMC acquired at 854 eV. (b and c) Euclidean distance map measured from the boundary and estimated Ni oxidation-state map of the thinned F-coated NMC secondary particle, respectively. The black shaded areas with white dotted lines in (b and c) indicate regions of the secondary particle covered with FIB welding material, which were excluded from the quantitative analysis. (d) Averaged XAS spectra across the O K-, F K-, and Ni L-edges for subgroups of pixels binned according to the distance map. The scatter points in the F K-edge XAS represent the measured values, and the smoothed curves are shown as guides to the eye. (e) OD difference map between the pre-edge (∼675 eV) and post-edge (∼710 eV) regions in the F K-edge spectral image stack. To reduce noise, six images from each region were averaged. The location of the zoomed difference map is indicated in (c) by the red dotted box. (f) Integrated F K-edge XAS intensity and averaged Ni oxidation state as functions of the distance from the secondary-particle surface. The example of the integration ranges is indicated in (d) by the gray shaded area. An example of linear-combination fitting using Ni²⁺ and Ni³⁺ reference spectra to reproduce the measured XAS is shown in Fig. S5 for the averaged whole-particle region and for single pixels marked by red crosses in (c).

Additionally, averaged XAS spectra across the O K-edge and Ni L-edge for subpixels binned by Euclidean distance from the secondary-particle surface (Fig. 2a and b) reveal the spatial distribution of Ni²⁺ and related surface species, as shown in Fig. 2d. The Ni²⁺ concentration in pristine F-coated NCM secondary particles is highest at the surface, decreasing rapidly within approximately 200 nm from the surface (Fig. 2d and f). Notably, the accuracy of the linear combination fitting of the reference spectra used to estimate the Ni²⁺ composition at each pixel was validated by the single-pixel spectral analysis results shown in Fig. S5b. The Ni oxidation state map (Fig. 2c), together with depth-sensitive Ni XAS analysis (Fig. 2d), reveals that the surfaces of the primary particles contain Ni²⁺ species and that the Ni²⁺ concentration increases progressively toward the surface of the secondary particles. The observed stoichiometric deviation likely arises from oxygen deficiency in the near-surface region of individual particles, which can promote local cation rearrangement and surface reconstruction under high-temperature calcination conditions,^62–66 and/or from the formation of a structurally disordered rock-salt- or spinel-like phase.¹⁰

The O K-edge absorption features, particularly in the low-energy region (<540 eV), are informative because of the pronounced sensitivity of this region to TM 3d–O 2p hybridized states.⁶⁷ Since three distinct peaks near 530, 532, and 535 eV correspond to Ni³⁺ 3d–O 2p hybridization, Ni²⁺ 3d–O 2p hybridization, and the characteristic absorption of surface by-products formed during high-temperature synthesis, such as Li₂CO₃, the depth-dependent spectral variation across the O K-edge, especially the attenuation of the 532 and 535 eV features relative to the peak near 530 eV as the distance from the surface increases, reveals a surface-localized distribution of a non-stoichiometric phase, represented as Li_1−xTM_1+xO₂ or LiMO_2−x, together with Li₂CO₃ by-products.^68,69 Additionally, residual lithium titration analysis showed that the total residual lithium content decreased after fluorination, although the Li₂CO₃ fraction increased slightly while the LiOH fraction decreased, as summarized in Table S1. Consistent with the observations in the STXM analysis, a minor amount of carbonate-containing surface species may form or remain during the PVDF-derived treatment.

The depth-dependent analysis of F K-edge spectra also confirms the localized presence of fluorine-containing species near the secondary-particle surface. Since the concentration of F species is relatively small, the averaged XAS across the whole particle region cannot reveal the existence of the F-containing species, as shown in Fig. S5a. In contrast, the OD difference between the pre-edge and post-edge regions in the F K-edge spectral images reveals clear contrast only near the surface of the secondary particle, indicating the localized presence of fluorine species, as shown in Fig. 2e. The depth-resolved XAS spectra and their smoothed integrated intensities, which are approximately proportional to the amount of F based on the assumption of homogenous chemical states, show that the F concentration in pristine F-coated NCM secondary particles is highest at the surface and decreasing rapidly within approximately 120 nm from the surface, as shown in Fig. 2d and f. Note that the estimated F coating thickness of approximately 120 nm may be overestimated because the long tails in the point spread function of an X-ray optic can cause spectral distortion near a boundary between two spectrally distinct regions.⁷⁰

To identify the chemical states of the fluorine species, additional STXM was utilized across the F K-edge (for more details, see the Experimental section and SI). Rapid survey scans based on two-energy difference mapping revealed that primary particles, obtained by fragmenting secondary particles, containing detectable fluorine species are exceedingly uncommon, with only about 26 instances found among over 200 particles examined (Fig. 3a and b). As shown in the two representative regions without detected fluorine species (Fig. 3a), the minor contrast observed arises from image misalignment caused by nonuniform particle motion during the STXM measurement. While the initial distribution of fluorine species within secondary particles prior to fragmentation cannot be conclusively determined, the present results support the hypothesis that fluorine is predominantly localized at the surfaces of secondary particles, with only a limited population detected within fragmented primary particles that likely originate from the secondary-particle surface. Possible origins of the fluorine species found on the surface of secondary NCM particles include residual PVDF remaining even after post-treatment, the formation of a LiF phase, or association with metal-fluorinated species (MF_x). LiF and PVDF can be discerned by their characteristic white-line features, appearing near 698 eV and approximately 3 eV below, respectively (Fig. 3c).^50,51,71 The averaged F K-edge spectrum acquired from the detected localized fluorine species, indicated by the blue dotted box in Fig. 3b, shows a main peak positioned 4 eV below the PVDF spectrum. This is a clear indication of the MF_x phase formation, as evidenced in the literature.^71,72

Fig. 3 Representative averaged optical density map (grayscale) and difference maps (color-coded from blue to red) obtained by comparing averaged pre-edge images (684 eV and 687 eV) with images acquired near the peak position (691 eV and 692 eV) for samples (a) without and (b) with detected fluorine species. (c) Averaged X-ray absorption spectroscopy (XAS, black line with scatter) extracted from the region containing fluorine species, as indicated by the blue dotted box in (b). For comparative purposes, reference spectra for PVDF (red solid line) and LiF (blue solid line) are also shown.

Based on the XPS, XRD, and STXM analyses, we can hypothesize the synthesis pathway of the fluorine-containing coating layer, as illustrated in Fig. 4. When the temperature is elevated in an inert atmosphere, the polymer structure of PVDF on the surface of the NCM cathode material decomposes, allowing the released fluorine ions to replace anions in the host NCM structure as well as leaving behind residual carbon on the surface. Since the potential presence of the LiF phase and PVDF residue was ruled out by XRD and STXM measurements, the coating layer is likely to exist in the form of a metal-fluoride. Although Ni–F bonding is most clearly resolved by XPS (Fig. 1f), the presence of minor Co–F and/or Mn–F species cannot be excluded. Therefore, the fluorinated surface phase is conservatively described here as MF_x-rich. As a result, the coating can protect oxygen release by the surface-localized metal fluoride phase and enhance electrochemical performance, as further demonstrated in the subsequent analysis.

Fig. 4 Schematic illustration of the PVDF-derived fluorination route and the resulting MF_x-rich surface layer on high-Ni NCM, shown together with residual NiO-like surface species.

Prior to evaluating the electrochemical performance under these harsh conditions, optimization studies were conducted to determine the appropriate PVDF concentration and annealing temperature. As summarized in Fig. S6 and S7, a PVDF content of 1 wt% and a heat treatment temperature of 350 °C were identified as optimal conditions, balancing initial capacity and cycling stability.

The electrochemical performances of the uncoated and F-coated NCMs were evaluated under harsh conditions, including an elevated operating-temperature of 40 °C and a high charge cut-off voltage of 4.5 V, to induce oxygen release followed by propagation of irreversible phases (Fig. 5). After the first two formation cycles at a rate of 20 mA g⁻¹, repeated charging steps were conducted at 100 mA g⁻¹ up to 4.5 V. The potential was held at 4.5 V until the current density dropped to 20 mA g⁻¹. The subsequent discharge step was carried out down to 2.7 V, also at a rate of 100 mA g⁻¹. The uncoated NCM initially exhibited a higher capacity; however, the F-coated NCM demonstrated superior capacity retention over repeated cycles (Fig. 5a and b). The higher overpotential observed during the initial charge cycle, along with the indistinct voltage plateau at higher voltages, suggests that a surface layer formed on the F-coated NCM impedes the movement of electrons and Li⁺. This is further supported by the lower specific discharge capacity in the first cycle, indicating that the surface layer initially hinders the electrochemical reaction but provides long-term benefits in terms of stability and capacity retention. Unlike the uncoated sample, the presence of an identifiable plateau in the high voltage regions at the beginning of discharge in the 50^th and 100^th cycles suggest that the surface layer formed on the coated sample acts as a passivation layer even under challenging conditions like high voltage and temperature (Fig. 5b).

Fig. 5 Voltage profiles of (a) uncoated and (b) F-coated NCM particles at the denoted cycles. (c) Cycle retention performance of both samples at 200 mA g⁻¹ current density. (d) Rate capabilities of both samples.

The enhanced electrochemical performance achieved through surface coating was confirmed by the discharge capacity retention up to the 250th cycle, as shown in Fig. 5c and Table 1. For the first formation cycle at a rate of 20 mA h g⁻¹, the uncoated and fluorine coated NCMs exhibited initial discharge capacities of 225.4 and 207.2 mA h g⁻¹, respectively. After increasing the rate to 200 mA g⁻¹ from the 3^rd cycle, however, the discharge capacity of the uncoated sample dropped significantly more (by approximately 23 mA h g⁻¹) compared to the coated sample (∼10 mA h g⁻¹). At a current density of 200 mA g⁻¹, the F-coated NCM exhibited capacity retention of approximately 75% and 70% at the 100^th and 150^th cycles, respectively, despite being subjected to high temperature and high voltage conditions. In contrast, the uncoated NCM showed lower capacity retention rates of only 65% and 58% over the same cycles, highlighting the improved stability and performance of the F-coated sample under harsh operating conditions. Furthermore, the enhanced stability imparted by the fluorine coating was demonstrated at room temperature (25 °C), as evidenced by the higher capacity retention of the F-coated NCM in SI Fig. S8. To ensure a fair comparison, additional capacity-matched cycling tests were performed by limiting both uncoated and F-coated NCM electrodes to the same charge capacity of 205 mA h g⁻¹ from the second cycle onward, and the F-coated NCM still showed slower capacity fading (Fig. S9), confirming the beneficial effect of the fluorine-derived surface treatment. The detailed experimental procedure is provided in the SI.

Table 1 Discharge capacity retention rate of uncoated NCM and F-coated NCM every 50^th cycles

Sample	50^th cycle	100^th cycle	150^th cycle	200^th cycle	250^th cycle
F-coated NCM	83.22%	74.20%	86.83%	64.88%	61.72%
Uncoated NCM	76.25%	65.07%	58.41%	53.97%	50.38%

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The difference in the capacity degradation rate at the elevated current rate aligns well with the expectation of suppression of oxygen release followed by formation of electrochemically inactive phases like NiO via surface fluorine coating.^5,73 The MF_x layer on the surface of the F-coated NCM, confirmed by XPS spectra and STXM, also retards thermal runaway owing to its stronger ion bonding. Consequently, owing to the passivation layer that prevents the generation of NiO, the F-coated NCM exhibited more stable cycling performance as well as higher rate capability, as shown in Fig. 5d. Moreover, upon returning from the high current density cycling to the low current density cycling, the capacity of the F-coated NCM remains stable compared to the initial capacity. In contrast to the F-coated NCM, the uncoated NCM exhibited more rapid capacity fading due to the formation an electrochemically inactive phase associated with sluggish Li⁺ transport. The battery performances of two NCM cathodes demonstrate that the fluorine-based surface layer hinders the formation and/or propagation of electrochemically less active phases like NiO.

The inhibition of NiO phase formation by the fluorine coating layer can be confirmed through the XPS spectra after 200 cycles within a voltage range of 2.7 to 4.5 V at a current density of 200 mA g⁻¹, as presented in Fig. 6a and b. When the charge and discharge cycles are repeated, particularly at high temperature and high voltage, the NiO rock-salt phase begins to grow from the surface inward into the cathode material, accompanied by oxygen release and accelerated irreversible structural changes.^5,11 As a result, oxygen loss related surface reconstruction during electrochemical cycling can be inferred from changes in the Ni XPS spectra. The intensity of Ni³⁺ in the uncoated NCM was initially dominant, but after cycling, the Ni²⁺ peak became more prominent (Fig. 6a). In contrast, the majority of the Ni³⁺ phase in the F-coated NCM remained largely unchanged even after 200 cycles. The shift in the ratio of Ni³⁺ to Ni²⁺ after cycling indicates that an effective passivation layer was formed on the surface of the F-coated NCM (Fig. 6b).

Fig. 6 Ni 2p XPS spectra for (a) uncoated and (b) F-coated samples after 200 cycles. Nyquist plots of uncoated NCM and F-coated NCM electrodes measured after (c) the 50th cycle and (d) the 100th cycle, illustrating the change in interfacial impedance during cycling. GITT-estimated Li⁺ diffusivity (e) as a function of voltage during the discharge step for the initial and 90^th cycles and (f) as a function of cycle number at 3.6 V.

Electrochemical impedance spectroscopy (EIS) and GITT were conducted to evaluate the evolution of interfacial transport behavior upon fluorination.^74–77 As shown in Fig. 6c and d, the Warburg impedance parameter σ (Ω sec^−1/2) and the corresponding apparent Li⁺ diffusivity (D_Li⁺, cm² s⁻¹), calculated from the Nyquist plots and GITT, show clear differences after 50 and 100 cycles as summarized in Table 2. In contrast to the trend observed after 50 cycles, clear divergences in both impedance and Li⁺ diffusivity emerged after prolonged cycling (100 cycles), demonstrating that the PVDF-derived fluorination suppresses diffusion-related impedance growth, consistent with improved interfacial stability under high-voltage operation and better preservation of reaction kinetics. Specifically, the diffusivity of Li⁺ in the uncoated NCM was equal to or higher than that of the F-coated NCM. The relatively higher lithium-ion diffusivity in the uncoated NCM is attributed to its relatively phase-pure surface and less cation mixing in the bulk structure, as gauged by the higher Ni³⁺ content in Ni 2p XPS spectra and the I₍₀₀₃₎/I₍₁₀₄₎ ratio in XRD, respectively. The absence of the NiO rock salt structure in both surface and bulk leads to the fast Li⁺ diffusion kinetics in the uncoated NCM at the beginning of cycling. However, after intensive electrochemical cycling, the Li⁺ diffusivity in the uncoated NCM decreased significantly compared to the F-coated sample. This is in line with predictions that the high-nickel cathode material will be resistant to oxygen release at high voltage due to the PVDF based metal-fluoride surface.⁷⁸

Table 2 Comparison of Warburg coefficients (σ) and Li⁺ diffusion coefficients (D_Li⁺) for uncoated and F-coated NCM electrodes after 50^th and 100^th cycles

σ	Uncoated NCM	F-coated NCM
50^th cycle	186.0435	187.9897
100^th cycle	248.2764	233.7336

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D Li^+	Uncoated NCM	F-coated NCM
50^th cycle	1.81 × 10^−16	1.77 × 10^−17
100^th cycle	1.02 × 10^−16	1.15 × 10^−16

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Taken together, the post-mortem XPS results, along with the EIS and GITT analyses, support a consistent structure–mechanism–performance relationship for the PVDF-derived fluorination strategy. After prolonged cycling, XPS indicates suppressed Ni²⁺-rich surface reconstruction in the F-coated NCM, while EIS shows mitigated interfacial impedance growth and GITT confirms better retention of Li⁺-transport kinetics during cycling. These coating-induced changes are consistent with the enhanced long-term cycling stability of the F-coated NCM under high-voltage conditions.

4. Conclusion

In this study, we developed a simple and scalable post-synthesis solid-phase treatment using PVDF to generate an ultrathin MF_x-rich surface layer on high-Ni LiNi_0.96Co_0.035Mn_0.005O₂ cathodes. Structural and spectroscopic analyses showed that the fluorine-derived interfacial layer formed without altering the bulk layered crystal structure. The modified surface suppressed NiO-like rock-salt formation and oxygen-release-related surface degradation, resulting in improved cycling stability, rate capability, and Li⁺ diffusivity under high-voltage (4.5 V) and elevated-temperature (40 °C) conditions. Because the proposed strategy is applied to already synthesized cathode particles through a simple and controllable process, it offers practical compatibility for stabilizing high-Ni cathodes in next-generation lithium-ion batteries.

Author contributions

HeeSang Lee: writing – original draft, visualization, validation, software, methodology, investigation, formal analysis, data curation, conceptualization. Wonchan Hwang: writing – original draft, visualization, software, methodology, investigation, formal analysis, data curation. Jahun Koo: writing – original draft, visualization, validation, software, methodology, investigation, formal analysis, data curation, conceptualization. Hendrik Ohldag: visualization, software, methodology, investigation, formal analysis, data curation. David A. Shapiro: visualization, software, methodology, investigation, formal analysis, data curation. Eun-Jung Shin: visualization, investigation, formal analysis, data curation. Min-Su Kim: visualization, investigation, formal analysis, data curation. Moonjung Jung: investigation, formal analysis, data curation. Namdong Kim: investigation, formal analysis, data curation. Jungjin Park: writing – review & editing, validation, supervision, project administration, funding acquisition, conceptualization. Young-Sang Yu: writing – review & editing, validation, supervision, project administration, funding acquisition, visualization, validation, software, investigation, formal analysis, data curation, conceptualization. Chunjoong Kim: writing – review & editing, validation, supervision, project administration, funding acquisition, conceptualization.

Conflicts of interest

There are no conflicts to declare.

Data availability

The data supporting this article have been included as part of the supplementary information (SI). Supplementary information: detailed procedures for precursor synthesis and fluorination, electrochemical testing including EIS, GITT, and capacity-matched cycling, STXM acquisition and analysis, as well as figures (Fig. S1–S10) and Table S1. See DOI: https://doi.org/10.1039/d6ta00410e.

Acknowledgements

This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korea Government (No. RS-2021-NR060128, RS-2024-00405185, RS-2024-00358571, and RS-2023-00284081) the National Research Council of Science & Technology (NST) grant (No. GTL24012-000), funded by the MSIT, Korea. This research was supported by the Regional Innovation System & Education (RISE) program through the (Chungbuk Regional Innovation System & Education Center), funded by the Ministry of Education (MOE) and the (Chungcheongbuk-do), Republic of Korea (2025-RISE-11-014-05). This research used resources of the Advanced Light Source, which is a DOE Office of Science User Facility under contract no. DE-AC02-05CH11231. J. P. was supported by the Institutional Program of the Korea Institute of Science and Technology (KIST) (No. 2E33941 and No. 2E3394B).

References

1. M. S. Whittingham, Chem. Rev., 2004, 104, 4271–4301.
2. J. B. Goodenough and Y. Kim, Chem. Mater., 2010, 22, 587–603.
3. W. Li, E. M. Erickson and A. Manthiram, Nat. Energy, 2020, 5, 26–34.
4. D. Mohanty, A. S. Sefat, S. Kalnaus, J. Li, R. A. Meisner, E. A. Payzant, D. P. Abraham, D. L. Wood and C. Daniel, J. Mater. Chem. A, 2013, 1, 6249–6261.
5. H. Zhang, H. Liu, L. F. J. Piper, M. S. Whittingham and G. Zhou, Chem. Rev., 2022, 122, 5641–5681.
6. A. Manthiram and W. Choi, Electrochem. Solid-State Lett., 2007, 10, A228.
7. E. McCalla, A. M. Abakumov, M. Saubanère, D. Foix, E. J. Berg, G. Rousse, M.-L. Doublet, D. Gonbeau, P. Novák, G. Van Tendeloo, R. Dominko and J.-M. Tarascon, Science, 2015, 350, 1516–1521.
8. M. Sathiya, A. M. Abakumov, D. Foix, G. Rousse, K. Ramesha, M. Saubanère, M. L. Doublet, H. Vezin, C. P. Laisa, A. S. Prakash, D. Gonbeau, G. VanTendeloo and J. M. Tarascon, Nat. Mater., 2015, 14, 230–238.
9. K. Min and E. Cho, Phys. Chem. Chem. Phys., 2018, 20, 9045–9052.
10. H. Liu, M. Wolfman, K. Karki, Y.-S. Yu, E. A. Stach, J. Cabana, K. W. Chapman and P. J. Chupas, Nano Lett., 2017, 17, 3452–3457.
11. K.-W. Nam, S.-M. Bak, E. Hu, X. Yu, Y. Zhou, X. Wang, L. Wu, Y. Zhu, K.-Y. Chung and X.-Q. Yang, Adv. Funct. Mater., 2013, 23, 1047–1063.
12. G. W. Nam, N.-Y. Park, K.-J. Park, J. Yang, J. Liu, C. S. Yoon and Y.-K. Sun, ACS Energy Lett., 2019, 4, 2995–3001.
13. J. Hong, H.-D. Lim, M. Lee, S.-W. Kim, H. Kim, S.-T. Oh, G.-C. Chung and K. Kang, Chem. Mater., 2012, 24, 2692–2697.
14. J.-Y. Luo, W.-J. Cui, P. He and Y.-Y. Xia, Nat. Chem., 2010, 2, 760–765.
15. S. Klein, P. Bärmann, T. Beuse, K. Borzutzki, J. E. Frerichs, J. Kasnatscheew, M. Winter and T. Placke, ChemSusChem, 2021, 14, 595–613.
16. Q. Wang, P. Ping, X. Zhao, G. Chu, J. Sun and C. Chen, J. Power Sources, 2012, 208, 210–224.
17. S. S. Zhang, Energy Storage Mater., 2020, 24, 247–254.
18. M. Jiang, D. L. Danilov, R.-A. Eichel and P. H. L. Notten, Adv. Energy Mater., 2021, 11, 2103005.
19. T. Hatsukade, A. Schiele, P. Hartmann, T. Brezesinski and J. Janek, ACS Appl. Mater. Interfaces, 2018, 10, 38892–38899.
20. J. Wandt, A. T. S. Freiberg, A. Ogrodnik and H. A. Gasteiger, Mater. Today, 2018, 21, 825–833.
21. R. Jung, M. Metzger, F. Maglia, C. Stinner and H. A. Gasteiger, J. Electrochem. Soc., 2017, 164, A1361–A1377.
22. J. Choi, J. Kim, K.-T. Lee, J. Lim, J. Lee and Y. S. Yun, Adv. Mater. Interfaces, 2016, 3, 1600784.
23. X. Xiong, Z. Wang, X. Yin, H. Guo and X. Li, Mater. Lett., 2013, 110, 4–9.
24. L. Sharma, M. Yi, E. Jo, H. Celio and A. Manthiram, Chem. Mater., 2022, 34, 4514–4522.
25. J. Wang, C. Du, C. Yan, X. Xu, X. He, G. Yin, P. Zuo, X. Cheng, Y. Ma and Y. Gao, RSC Adv., 2016, 6, 26307–26316.
26. K.-J. Park, M.-J. Choi, F. Maglia, S.-J. Kim, K.-H. Kim, C. S. Yoon and Y.-K. Sun, Adv. Energy Mater., 2018, 8, 1703612.
27. B.-B. Lim, S.-J. Yoon, K.-J. Park, C. S. Yoon, S.-J. Kim, J. J. Lee and Y.-K. Sun, Adv. Funct. Mater., 2015, 25, 4673–4680.
28. T. Weigel, F. Schipper, E. M. Erickson, F. A. Susai, B. Markovsky and D. Aurbach, ACS Energy Lett., 2019, 4, 508–516.
29. H. Kim, S.-B. Kim, D.-H. Park and K.-W. Park, Energies, 2020, 13, 4808.
30. U.-H. Kim, G.-T. Park, B.-K. Son, G. W. Nam, J. Liu, L.-Y. Kuo, P. Kaghazchi, C. S. Yoon and Y.-K. Sun, Nat. Energy, 2020, 5, 860–869.
31. J. Yan, H. Huang, J. Tong, W. Li, X. Liu, H. Zhang, H. Huang and W. Zhou, Interdiscip. Mater., 2022, 1, 330–353.
32. Z. Chen, Y. Qin, K. Amine and Y. K. Sun, J. Mater. Chem., 2010, 20, 7606–7612.
33. Y. K. Sun, S. W. Cho, S. W. Lee, C. S. Yoon and K. Amine, J. Electrochem. Soc., 2007, 154, A168.
34. Y.-K. Sun, M.-J. Lee, C. S. Yoon, J. Hassoun, K. Amine and B. Scrosati, Adv. Mater., 2012, 24, 1276.
35. C.-D. Li, J. Xu, J.-S. Xia, W. Liu, X. Xiong and Z.-A. Zheng, Solid State Ionics, 2016, 292, 75–82.
36. S. Dai, G. Yan, L. Wang, L. Luo, Y. Li, Y. Yang, H. Liu, Y. Liu and M. Yuan, J. Electroanal. Chem., 2019, 847, 113197.
37. J. Ahn, E. K. Jang, S. Yoon, S.-J. Lee, S.-J. Sung, D.-H. Kim and K. Y. Cho, Appl. Surf. Sci., 2019, 484, 701–709.
38. Y. Bi, W. Yang, R. Du, J. Zhou, M. Liu, Y. Liu and D. Wang, J. Power Sources, 2015, 283, 211–218.
39. Y.-S. He, L. Pei, X.-Z. Liao and Z.-F. Ma, J. Fluorine Chem., 2007, 128, 139–143.
40. A. R. Naghash and J. Y. Lee, Electrochim. Acta, 2001, 46, 2293–2304.
41. Q. Shi, F. Wu, H. Wang, Y. Lu, J. Dong, J. Zhao, Y. Guan, B. Zhang, R. Tang, Y. Liu, J. Liu, Y. Su and L. Chen, Energy Storage Mater., 2024, 67, 103264.
42. Q. Dong, J. Wu, Y. Wang, J. Qi, J. Wang, J. Li, L. Zhao, S. Hu, H. Shao, Y. Shen and L. Chen, Energy Storage Mater., 2023, 63, 103054.
43. H. Ding, Y. Su, X. Wang, Y. Hu, X. Li, H. Zhang, G. Liu, W. Yu, X. Dong, J. Wang and X. Wang, J. Energy Chem., 2025, 108, 427–457.
44. I. Belharouak, Y. K. Sun, J. Liu and K. Amine, J. Power Sources, 2003, 123, 247–252.
45. C. A. Schneider, W. S. Rasband and K. W. Eliceiri, Nat. Methods, 2012, 9, 671–675.
46. T. Feggeler, A. Levitan, M. A. Marcus, H. Ohldag and D. A. Shapiro, J. Electron Spectrosc. Relat. Phenom., 2023, 267, 147381.
47. P. M. Csernica, S. S. Kalirai, W. E. Gent, K. Lim, Y.-S. Yu, Y. Liu, S.-J. Ahn, E. Kaeli, X. Xu, K. H. Stone, A. F. Marshall, R. Sinclair, D. A. Shapiro, M. F. Toney and W. C. Chueh, Nat. Energy, 2021, 6, 642–652.
48. B. L. Henke, E. M. Gullikson and J. C. Davis, At. Data Nucl. Data Tables, 1993, 54, 181–342.
49. Y.-S. Yu, M. Farmand, C. Kim, Y. Liu, C. P. Grey, F. C. Strobridge, T. Tyliszczak, R. Celestre, P. Denes, J. Joseph, H. Krishnan, F. R. N. C. Maia, A. L. D. Kilcoyne, S. Marchesini, T. P. C. Leite, T. Warwick, H. Padmore, J. Cabana and D. A. Shapiro, Nat. Commun., 2018, 9, 921.
50. K. Y. Chung, W.-S. Yoon, K.-B. Kim, B.-W. Cho and X.-Q. Yang, J. Appl. Electrochem., 2011, 41, 1295–1299.
51. M. Sina, R. Thorpe, S. Rangan, N. Pereira, R. A. Bartynski, G. G. Amatucci and F. Cosandey, J. Phys. Chem. C, 2015, 119, 9762–9773.
52. H.-J. Shin, N. Kim, H.-S. Kim, W.-W. Lee, C.-S. Lee and B. Kim, J. Synchrotron Radiat., 2018, 25, 878–884.
53. X. Li, Z. Xie, W. Liu, W. Ge, H. Wang and M. Qu, Electrochim. Acta, 2015, 174, 1122–1130.
54. J. Morales, C. Pérez-Vicente and J. L. Tirado, Mater. Res. Bull., 1990, 25, 623–630.
55. T. Ohzuku, A. Ueda and M. Nagayama, J. Electrochem. Soc., 1993, 140, 1862.
56. R. V. Moshtev, P. Zlatilova, V. Manev and A. Sato, J. Power Sources, 1995, 54, 329–333.
57. J. U. Choi, N. Voronina, Y.-K. Sun and S.-T. Myung, Adv. Energy Mater., 2020, 10, 2002027.
58. S. H. Lee, S. Lee, B. S. Jin and H. S. Kim, Sci. Rep., 2019, 9, 8901.
59. Y. Zong, Z. Guo, T. Xu, C. Liu, Y. Li and G. Yang, Int. J. Energy Res., 2020, 44, 8532–8541.
60. N. Zhang, J. Stark, H. Li, A. Liu, Y. Li, I. Hamam and J. R. Dahn, J. Electrochem. Soc., 2020, 167, 080518.
61. C. J. Powell and A. Jablonski, Nucl. Instrum. Methods Phys. Res., Sect. A, 2009, 601, 54–65.
62. S. Levasseur, M. Ménétrier, Y. Shao-Horn, L. Gautier, A. Audemer, G. Demazeau, A. Largeteau and C. Delmas, Chem. Mater., 2003, 15, 348–354.
63. J.-M. Lim, D. Kim, Y.-G. Lim, M.-S. Park, Y.-J. Kim, M. Cho and K. Cho, ChemElectroChem, 2016, 3, 943–949.
64. A. R. Armstrong, M. Holzapfel, P. Novák, C. S. Johnson, S.-H. Kang, M. M. Thackeray and P. G. Bruce, J. Am. Chem. Soc., 2006, 128, 8694–8698.
65. Y. Xia, H. Wang, Q. Zhang, H. Nakamura, H. Noguchi and M. Yoshio, J. Power Sources, 2007, 166, 485–491.
66. Z.-K. Tang, Y.-F. Xue, G. Teobaldi and L.-M. Liu, Nanoscale Horiz., 2020, 5, 1453–1466.
67. F. M. F. de Groot, M. Grioni, J. C. Fuggle, J. Ghijsen, G. A. Sawatzky and H. Petersen, Phys. Rev. B:Condens. Matter Mater. Phys., 1989, 40, 5715–5723.
68. C. Tian, D. Nordlund, H. L. Xin, Y. Xu, Y. Liu, D. Sokaras, F. Lin and M. M. Doeff, J. Electrochem. Soc., 2018, 165, A696.
69. W.-S. Yoon, O. Haas, S. Muhammad, H. Kim, W. Lee, D. Kim, D. A. Fischer, C. Jaye, X.-Q. Yang, M. Balasubramanian and K.-W. Nam, Sci. Rep., 2014, 4, 6827.
70. M. A. Marcus, D. A. Shapiro and Y.-S. Yu, Microsc. Microanal., 2021, 27, 1448–1453.
71. R. Qiao, I. T. Lucas, A. Karim, J. Syzdek, X. Liu, W. Chen, K. Persson, R. Kostecki and W. Yang, Adv. Mater. Interfaces, 2014, 1, 1300115.
72. D. Leanza, C. A. F. Vaz, G. Melinte, X. Mu, P. Novák and M. El Kazzi, ACS Appl. Mater. Interfaces, 2019, 11, 6054–6065.
73. L. Zou, W. Zhao, Z. Liu, H. Jia, J. Zheng, G. Wang, Y. Yang, J.-G. Zhang and C. Wang, ACS Energy Lett., 2018, 3, 2433–2440.
74. S. D. Talian, S. Brutti, M. A. Navarra, J. Moškon and M. Gaberscek, Energy Storage Mater., 2024, 69, 103413.
75. M. J. Christian and M. Alain, AIMS Mater. Sci., 2019, 6, 406–440.
76. D. W. Dees, S. Kawauchi, D. P. Abraham and J. Prakash, J. Power Sources, 2009, 189, 263–268.
77. A. Nickol, T. Schied, C. Heubner, M. Schneider, A. Michaelis, M. Bobeth and G. Cuniberti, J. Electrochem. Soc., 2020, 167, 090546.
78. F. T. Geldasa, M. A. Kebede, M. W. Shura and F. G. Hone, RSC Adv., 2022, 12, 5891–5909.

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Footnote

† H. L., W. H., and J. K contributed equally to this work.

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