Kinetically Dormant Ni‐Rich Layered Cathode During High‐Voltage Operation

Abstract

The degradation of Ni‐rich cathodes during long‐term operation at high voltage has garnered significant attention from both academia and industry. Despite many post‐mortem qualitative structural analyses, precise quantification of their individual and coupling contributions to the overall capacity degradation remains challenging. Here, by leveraging multiscale synchrotron X‐ray probes, electron microscopy, and post‐galvanostatic intermittent titration technique, the thermodynamically irreversible and kinetically reversible capacity loss is successfully deconvoluted in a polycrystalline LiNi_0.83Mn_0.1Co_0.07O₂ cathode during long‐term charge/discharge cycling in full cell configuration. Contradicting the dramatic capacity loss, the layered structure remains highly alive even after 1000 cycles at 4.6 V while undergoing a three‐order of magnitude reduction in the mass transfer kinetics, leading to almost fully recoverable capacity under kinetic‐free conditions. Such kinetic dormant behavior after cycling is not simply ascribed to poor chemical diffusion by reconstructed cathode surface but highly synchronizes with the lattice strain evolution stemming from the structural heterogeneity between deeply delithiated layered and degraded rock‐salt phases at high voltage. These findings deepen the degradation mechanism of high‐voltage cathodes to achieve long‐cycling and fast‐charging performance.

Ni‐rich layered cathode remains thermodynamically reversible but undergoes dramatic kinetic degradation at high voltage, leading to the commonly observed capacity loss that in fact can be recovered under kinetic‐free conditions. Such kinetic dormant behavior highly synchronizes with the lattice strain evolution stemming from layered/rock‐salt structural heterogeneity during charge.

1. Introduction

Layered LiNi_1‐x‐yMn_xCo_yO₂ (NMC) cathode materials are among the most promising candidates for next‐generation lithium‐ion and lithium‐metal batteries due to their high specific capacities.^{[ 1 ]} Pushing NMC cathodes to a higher voltage and higher Ni content is essential to unlocking high‐energy‐density batteries to meet the ever‐increasing demand for automotive applications. However, Ni‐rich cathodes operated at a relatively high potential (>4.3 V vs Li/Li⁺) suffer from a universal performance degradation during cycling.^{[ 2 ]} Despite well‐demonstrated improvements by various trial‐and‐error strategies including surface coating,^{[ 3 ]} element doping,^{[ 4 ]} and electrolyte engineering,^{[ 5 ]} the degradation mechanism of Ni‐rich cathodes at high voltage remains controversial, which strongly limits their wide applications toward extreme conditions.

Early study of surface reconstruction and chemical evolution of layered cathodes revealed a layered‐to‐spinel/rock‐salt phase transformation and formation of surface reaction layer after a high voltage charge.^{[ 2d ]} The reconstructed surface significantly blocks the diffusion of lithium ions and electrons and increases the cell impedance, resulting in severe intra/inter‐particle heterogeneity.^{[ 6 ]} On the other hand, layered cathodes undergo an inherent anisotropic lattice variation at deep charge irrespective of Ni content.^{[ 2a ]} The large lattice mismatch between the bulk layered structure and surface rock‐salt phase created high interfacial lattice strain and eventually caused irreversible structural degradation of bulk fatigue at high voltage after repeated cycling.^{[ 7 ]} As a result, chemo‐mechanical cracking has also been often observed in the cycled Ni‐rich cathode particles,^{[ 2} , ^{8 ]} which is believed to aggravate the side reactions with electrolytes and induce the formation of thick ionic‐isolating cathode‐electrolyte interphase (CEI). Furthermore, the adverse effects of transition metal dissolution, migration, and deposition on the anode have been often discussed.^{[ 9 ]} In addition to this cycling‐initiated structural degradation, the synthesis‐induced native structural defects such as tilted boundaries^{[ 10 ]} and dislocations^{[ 11 ]} also contributed to the premature capacity degradation of Ni‐rich cathodes during cycling.

Although these valuable findings are acknowledged, no consensus has been reached on the dominant degradation pathway of Ni‐rich cathodes. For example, Manthiram et al. argued that particle cracking is not the main cause of the capacity fade of Ni‐rich cathodes but rather a symptom of increased surface reactivity.^{[ 12 ]} The degradation of single‐crystalline cathodes without grain boundaries at high voltage also challenges the prevailing belief of intergranular cracking on capacity loss.^{[ 13 ]} Surprisingly, a comparison between single‐crystalline and polycrystalline cathodes shows that cracking of secondary particles in polycrystalline cathodes decreases the charge‐transfer resistance and improves the Li⁺ diffusion coefficient due to electrolyte infiltration, boosting reversible capacity and rate capability.^{[ 14 ]} These conflict statements based on qualitative structural analysis lies in lacking a precise quantification of their individual and coupling contributions to the overall capacity loss of Ni‐rich cathodes. What is worse, most of these observations are deduced from half‐cell tests with lithium metal as a reference and counter electrode. However, the high reactivity of lithium metal with organic electrolytes contributes significantly to cell degradation, complicating an accurate analysis of the capacity loss of the cathode itself.^{[ 15 ]}

In this work, it is our major intent to reveal the dominant degradation pathway of Ni‐rich cathodes during long‐term operation at high voltage. Specifically, we used a conventional polycrystalline LiNi_0.83Mn_0.1Co_0.07O₂ (Ni83) cathode and evaluated its long‐term cycling performance using full cell configuration with commercial graphite anode to exclude potential interferences from lithium metal anode. In addition to the use of a suit of synchrotron X‐ray probes and electron microscopy to comprehensively probe surface and bulk structural evolution, we employed a post‐galvanostatic intermittent titration technique (GITT) method to probe the harvested cathodes from full cells, which includes plenty discrete pulses (>300 steps in lithiation/delithiation) and the relaxation potential for state‐of‐charge (SOC) reference, to better determine both thermodynamic equilibrium behavior and kinetic response of the subtle equilibrium changes. Contradicting the commonly observed capacity fading after high‐voltage cycling, gathered experimental evidence collectively shows that the layered structure of Ni83 cathode remains highly active even after 1000 cycles at 4.6 V but undergoes a dramatic kinetic degradation by three orders of magnitude. Such kinetically limited capacity loss can be almost fully recovered under kinetic‐free conditions, which runs counter to the prevailing perception of thermodynamically irreversible capacity loss by bulk structural degradation. Moreover, the aggravated kinetic dormant behavior at high voltage is not solely driven by the chemical diffusion reduction of ionic‐isolating CEI due to the exacerbated side reactions but highly synchronizes with the lattice strain evolution stemming from layered/rock‐salt structural heterogeneity at deep delithiation state. Through deconvolution of thermodynamic irreversible and kinetic reversible capacity loss, this work resolves the long‐standing controversial understanding of the degradation mechanism of Ni‐rich cathodes at high‐voltage and paves the way to developing high‐voltage cathodes for high‐energy and long‐life batteries.

2. Results

2.1. Electrochemical Performance Evaluation and Post‐Mortem Material Diagnosis

We selected a Ni83 cathode as the model material due to its good quality in particle morphology and crystal structure (Figure S1, Supporting Information). Considering that the chemically unstable metallic lithium complicates the interpretation of the electrochemical evaluation,^{[ 15 ]} the long‐term cycling performance of Ni83 cathode at various SOCs (i.e., 4.2, 4.4, and 4.6 V) was evaluated in the practical full‐cell configuration with the graphite anode, as shown in Figure 1a and Table S1 (Supporting Information). The detailed preparation of coupled full cells with targeted voltage windows and N/P capacity ratio is discussed in Figure S2 (Supporting Information) of Supporting Information. Ni83 cathode at 4.2 V in full cells exhibited 89.3% capacity retention (Figure 1a) and 99.72% average coulombic efficiency (CE in Figure S3, Supporting Information) after long‐term cycling at C/10, and the discharge mean voltage of the cell was 3.68 V (Figure S4a, Supporting Information), almost identical to the initial value. The complete recovery of the reversible capacity and the discharge mean voltage at C/10 suggests the interfacial impedance hike as dominated failure mechanisms, with minor contribution from either cathode or anode material degradation. In comparison, the Ni83 cathode at 4.4 V in full cells exhibited a lower capacity retention of 74.2% and an average CE of 99.69% (Figure S3, Supporting Information) after cycling, and the discharge mean voltage rapidly decayed from 3.71 V at 1st cycle to 3.59 V at 1010th cycle (Figure S4b, Supporting Information). The large capacity loss and unrecovered discharge mean voltage indicate increased deterioration from electrode materials. As of our greatest interest in mechanistic investigation, the extended practice was conducted on Ni83 cathode at a higher cut‐off voltage of 4.6 V in full cells to exacerbate the degradation. Full cells exhibited only 40.0% retention and 99.03% average CE (Figure S3, Supporting Information) with a dramatic reduction in the discharge mean voltage (Figure S4c, Supporting Information).

Figure 1.

Long‐Term Cycling of the Model Full Cells and the Observed Issues on the Cathode and Anode as Commonly Reported. a) Electrochemical evaluations of Ni83 at 4.2, 4.4, and 4.6 V in full cells against graphite anode for two formation cycles at C/10, 1000 testing cycles at 1C, and 10 post‐test cycles at C/10. High‐resolution STEM images and selected area FFT patterns compare the local surface structure on b) pristine Ni83 cathode and c) aged Ni83 cathode after 1000 cycles at 4.6 V, showing the exacerbated surface reconstruction of rock‐salt structure during high voltage operation. The comparison of surface and interior structure between pristine material and cycled cathodes at different voltages is shown in Figure S5 (Supporting Information). e) Soft XAS Ni L‐edge spectra in TEY and TFY channels for surface and bulk information in pristine and cycled cathodes at various upper voltages, respectively. The reference spectra of Ni²⁺, Ni³⁺, and Ni⁴⁺ are provided to reveal the Ni valence state change at the utmost surface. XRF color maps of graphite electrode surfaces with quantitative analysis of migrated transition metals in the full cells coupled with Ni83 cathode at f) 4.2 V, g) 4.4 V, and h) 4.6 V.

Post‐mortem diagnosis was performed on the harvested Ni83 cathodes and graphite anodes from full cells after cycling at different upper voltages. Through various surface characterizations, the cycled cathodes showed more severe surface reconstructions at elevated charge voltages, such as the undesired rock‐salt surface layer (HR‐STEM in Figure 1b,c and TEM in Figure S5, Supporting Information) and more surface Ni reduction (i.e., Ni³⁺/Ni⁴⁺ to Ni²⁺) via Ni L‐edge soft XAS (Figure 1d). It is prevailingly believed that the reconstructed cathode surface substantially impedes the chemical diffusion of lithium ions to build up the cell impedance and results in severe reaction heterogeneity within inter/intra‐particles for potential cracking. In addition, the dissolution of transition metals due to parasitic reactions of deep‐charged cathodes with electrolytes can lead to active material loss and irreversible capacity loss. By leveraging the quantitative measurements of synchrotron X‐ray fluorescence (XRF), the amount of migrated transition metals at the anode increased dramatically with elevated charge voltage, like 73.29 vs 0.97 µg cm⁻² Ni content with Ni83 cathode at 4.6 and 4.2 V, respectively (Figure 1e–g). Notably, it is found that the total amount of migrated and dissolved transition metals (including HPIC measurements in the collected electrolyte in Table S2, Figure S6, Supporting Information) is negligible like ≈110 µg cm⁻² Ni with Ni83 cathode at 4.6 V, compared to the active cathode loading of ≈10 mg cm⁻². However, the dissolution of transition metals would practically bring a more detrimental crosstalk influence on graphite anode for the electrocatalytic reduction to accelerate the electrolyte decomposition, SEI growth (Figure S7, Supporting Information), and capacity degradation (Figure S8, Supporting Information). These structural analysis results well align with previous findings of cycled Ni‐rich cathodes at high voltage and have been qualitatively correlated with the observed capacity loss.^{[ 2} , ⁶ , ^{9 ]} However, quantification of their individual and coupling contributions remains challenging.

To further deconvolute the capacity loss, full cells after long‐term cycling were fully discharged at C/10 and then dissembled, and new half cells using the harvested electrodes were assembled with fresh lithium metal anode and electrolyte. The recovered Ni83 cathode in half cells exhibited 161 mAh g⁻¹ and 86% capacity retention at 4.2 V (Figure 2a), 160 mAh g⁻¹ and 72% retention at 4.4 V (Figure 2b), and 141 mAh g⁻¹ and 60% retention at 4.6 V (Figure 2c) at C/10 rate. Surprisingly, the bulk material characterization via synchrotron high‐energy X‐ray diffraction (HE‐XRD) revealed that cycled Ni83 cathodes at all upper cutoff voltages still retained the layered structure similar to pristine material, without significant degradation to spinel or rock‐salt structure (Figure 2d and Table S3, Supporting Information). Likewise, the transmission electron microscope (TEM, Figure S5a–d, Supporting Information) showed a well‐ordered layered structure at the interior of the cathode materials as the pristine material. The X‐ray absorption fine structure (XAFS) spectra also reflected insignificant variations in the valence states and the neighboring atomic coordination of transition metals in the bulk material (Figure 2e,f; Figure S9, Supporting Information). Moreover, the widely concerned intergranular cracking, due to the abrupt anisotropic lattice change in the secondary particles at high voltage,^{[ 16 ]} was also absent via the cross‐section SEM. Compared with the compacted particles in the pristine cathode, subtle cracks, and voids were observed within secondary particles after long‐term cycling at high voltage. Nevertheless, no significant particle pulverization was observed at all potentials after 1000 cycles, especially in the large scanning area of multiple secondary particles cycled at 4.6 V (Figure 2g and see the morphology of pristine and cycled particles at 4.2 and 4.4 V in Figure S10, Supporting Information). These observations conflict with the severe capacity loss after cycling, challenging the conventional perception of several structural degradations for capacity loss.

Figure 2.

Inconsistency between Electrochemical Evaluation and Bulk‐Type Post‐Mortem Diagnoses of the Harvested Cathodes in Full Cells. The discharging profiles of full cells using Ni83 cathodes at a) 4.2 V, b) 4.4 V, and c) 4.6 V in the formation cycle at C/10, long‐term cycling of 5th, 500th, and 1000th cycles at 1C and post‐test cycle at C/10, as well as the half cells of pristine and harvested cathodes at C/10. The cathode at higher voltage suffers more severe capacity fading. d) Synchrotron‐based ex situ HE‐XRD characterizations exhibit good, layered structure without significant rock‐salt or spinel phase in all cycled Ni83 cathodes at various upper voltages after 1000 cycles, compared with pristine Ni83 cathode. The X‐ray wavelength is 0.1173 Å. e) XANES and f) EXAFS spectra of Ni K‐edge to reveal the slight evolutions of valence states and radical distances in all cycled Ni83 cathodes at various upper voltages. The Co and Mn spectra are shown in Figure S9 (Supporting Information). g) The cross‐sectional SEM images in a large view of multiple Ni83 particles at 4.6 V after 1000 cycles (see pristine and cycled cathodes at 4.2 and 4.4 V in Figure S10, Supporting Information). Inlet: zoom‐in view showing some narrow intergranular cracks throughout the entire particle. No significant particle fracture for the isolation of active material is seen at 4.6 V after 1000 cycles.

2.2. Performance Deterioration Deconvolution and Key Insights from Kinetic Evaluation

Besides thermodynamic redox reactions of active materials, the electrochemistry of batteries also critically depends on the mass‐transfer kinetics from surface to bulk in the electrode material. Heretofore, few studies have quantitatively evaluated the deterioration contributions of thermodynamic irreversible capacity loss and/or kinetic reversible capacity loss in Ni‐rich cathode after long‐term cycling at various SOCs. Therefore, deconvoluting the overall performance deterioration into the irreversible loss and the reversible loss is particularly targeted in this study to seek the dominant deterioration mechanism.

To deconvolute the contributions of material degradation and kinetic limit, we have carried out post‐GITT tests on the harvested cathodes (after 1000 cycles in full cells) with half cells using fresh Li metal anode and fresh electrolyte. In order to reach the thermodynamic equilibrium state in the cathode material with a minimum concentration gradient, our post‐GITT protocol employed plenty of discrete pulses (>300 steps in lithiation/delithiation) and the relaxation potential for SOC reference, which help determine both thermodynamic equilibrium behavior and kinetic response of the subtle equilibrium changes. The practical capacity of alive cathode material can be measured from the discharge steps of galvanostatic current pulses by minimizing the influence of kinetic limits. Remarkably, GITT results revealed that the cycled Ni83 cathodes after 1000 cycles at 4.2, 4.4, and 4.6 V still sustained ≈96%, 92%, and 89% capacity retention under the “kinetic‐free” circumstances (Figure 3a), respectively. Consistent with insignificant bulk material degradation via post‐mortem characterizations, the irreversible capacity loss due to the active material loss is minor in the overall capacity loss of Ni‐rich cathodes. Thus, the significant capacity difference between GITT tests and the constant‐current tests in half cells (Figure 2a–c) represents the contribution of kinetic limits.

Figure 3.

Performance Deconvolution to Unveil the Critical Contribution of Kinetic Limits. a) GITT profiles of Ni83ǀǀLi half cells using pristine Ni83 cathode and the cycled Ni83 cathodes after 1000 cycles in full cells in different voltage windows, as well as the comparison of their discharge capacity. Inlet images show their single iteration of the current pulse and relaxation processes. b) A quantitative reference of the deterioration contribution of Li reservoir loss, kinetic limits, or material degradation, individually, in Ni83 cathodes in different voltage windows after 1000 cycles. The contribution from kinetic limits is defined as the capacity gap between the constant‐current tests and the GITT tests of the same harvested cathode half cells. The contribution from material degradation is determined by the capacity gap in GITT tests between the harvested cathode and the pristine cathode. Li reservoir loss is assumed as the capacity gap between the 1010th‐cycle full cells and the first‐cycle harvested cathode half cells at the same current rate of C/10, representing the Li vacancy existing in cathode material for the incomplete re‐lithiation at the lower end of SOC.

The mass transfer kinetics, roughly regarded for now as the overall diffusion of Li (S²D_Li+ in GITT), and the overall electrical impedance (I^*R drop in GITT) were taken into account for investigating the kinetic limit mechanism. Since fresh electrolyte and Li metal anode were used in the recovered half cells, the measured diffusion and I^*R drop via GITT tests can be good indicators of the kinetic properties of the Ni83 cathode. Because the active surface area is difficult to accurately measure in porous electrodes, the surface‐independence chemical diffusion (S²D_Li+ ) was determined for all Ni83 cathodes. The measured S²D_Li+ values of ≈5E‐8 cm⁶ s⁻¹ in the pristine NMC cathode material show the same magnitude as the reported literature^{[ 17 ]} (be aware of different cathode surface areas). Compared with the intrinsically almost homogeneous kinetic properties in the pristine Ni83 cathode, the cycled cathode exhibited a dramatic reduction in mass transfer kinetics as a function of upper potentials, showing almost three‐order lower magnitude at 4.6 V. Yet, the over‐voltage gain (I^*R) or the overall electrical impedance was relatively insignificant in the kinetic influence of cycled Ni83 cathode, such as increasing 1.5 times at 4.2 V to 4.1 times at 4.6 V. Hence, the reduction in the mass‐transfer kinetics is of higher significance toward the kinetic deteriorations during long‐term cycling of Ni83 cathodes. To shed light on the contributions of different deterioration mechanisms after long‐term cycling, Figure 3b presents a quantitative reference of the individual capacity contribution from material degradation, kinetic limits, and Li reservoir loss at different upper cutoff potentials. The comparison reveals that the kinetic limit is the predominant deterioration mechanism with the most capacity loss, like 29% loss at 4.6 V, while the overall material degradation directly leads to relatively less irreversible deterioration during cycling of Ni83 cathode, like 11% loss at 4.6 V.

The cycling evolution of Ni83 cathodes in full cells at the 1st, 100th, 300th, 500th, and 1000th cycles with different charge cut‐off voltages were further investigated to probe more insights about the impeded mass‐transfer kinetics (see cycling tests at 4.6V in Figure 4a and other voltages in Figures S11–S13, Supporting Information). Figure 4b exhibits the cycling evolutions of kinetic reversible loss and thermodynamically irreversible loss of Ni83 cathodes at different charge cut‐off voltages, revealing predominant kinetic deterioration in all cycles and voltages. Likewise, the discharge mean voltage as the reaction characteristics can semi‐quantitatively reflect the minor thermodynamic material degradation from the subtle reduction in post‐GITT tests, while the dramatic reduction in the regular constant‐current tests confirms predominant kinetic deterioration (see Figure S14, Supporting Information). In addition, it is found that the kinetic reversible loss is highly modulated by the inhibited mass‐transfer kinetics at the H2‐H3 phase transition in the deep delithiation process (Figure 4c), suggesting the kinetically dormant H2‐H3 transition. Before further kinetic analysis, one should also know that the mass‐transfer kinetics in the electrode material is typically regarded as the chemical diffusion of Li⁺, but it is found in this study that the mass‐transfer kinetics does not exclusively depend on chemical diffusion as explained in the following session. But for now, the overall diffusion is simply used to elaborate the deteriorating evolution of mass‐transfer kinetics over cycling.

Figure 4.

Key Insights from Mass‐Transfer Kinetic Characteristics toward Phase Transitions in Cycled Ni‐Rich Cathode. a) Electrochemical cycling of Ni83 cathode at 4.6 V in full cells for 100, 300, 500, and 1000 cycles at 1C, followed by 3 cycles at C/10. b) Cycling‐evolutions of the quantitative kinetic reversible loss and thermodynamically irreversible loss in harvested Ni83 cathodes at three upper voltages. c) The reversible capacity losses in all harvested Ni83 cathodes show a strong dependence on the kinetic deterioration at H2‐H3 transition at high voltage. The 1st, 100th, 300th, 500th, and 1000th‐cycle evolutions of d–f) voltage‐dependent chemical diffusion of Li⁺, g–i)thermodynamic “equilibrium” phase transitions derived from GITT relaxation steps under “kinetic‐free” circumstance, and j–l) kinetic‐limited phase transitions from galvanostatic cycling tests at C/10 in Ni83 cathodes at a) 4.2 V (see raw data in Figure S11, Supporting Information), b) 4.4 V (Figure S12, Supporting Information), and c) 4.6 V, respectively. In diffusion plots, solid symbols represent the charging process and hollow symbols for the discharging process.

The cycling evolutions of the mass transfer kinetics (Figure 4d–f) reveal severe reduction over longer cycling and at higher voltage. More notably, the Ni83 cathodes cycled at high voltage exhibited inhomogeneous kinetic reduction with a dramatic drop at the deep delithiation process above 4.1 V, where the H2‐H3 phase transition occurs. One should note that the intrinsic kinetic properties of the H2‐H3 transition in the pristine cathode at 1st cycle are almost identical to other phase transitions at the lower voltage region. To elucidate the impact of this inhomogeneous kinetics on cathode performance, we demonstrate a “kinetic‐free” dQdV plot in this study using the end voltage values in GITT relaxation steps (later denoted as GITT‐dQdV) for the thermodynamically “equilibrium” phase transitions without kinetic limit (Figure 4g–i). GITT‐dQdV unearths that the prevailingly degraded H2‐H3 transition at high SOC region can be restored with highly thermodynamic reversibility under “kinetic‐free” circumstances, even at 4.6 V after 1000 cycles. In sharp contrast, the diminished redox reaction in the H2‐H3 transition was observed in the conventional dQdV plot from regular galvanostatic (or constant‐current) cycling tests at C/10 (Figure 4j–l), because of the substantial kinetic reduction at high voltage for predominant reversible capacity loss (Figure 4c). These findings confirm that Ni‐rich cathodes at high voltage suffer the kinetic‐dominant performance deterioration, against the prevailingly believed structural degradation mechanism of either lattice collapse or bulk fatigue structure that loses electrochemical activity of H2‐H3 phase transition. Despite its critical influence in high‐voltage performance deterioration, the origin of the aggravated kinetic reduction at the H2‐H3 transition is barely explored by far. This inhomogeneous behavior cannot be simply interpreted by current knowledge of the ionic‐isolating surface reconstruction layer^{[ 18 ]} and the concentration‐dependent chemical diffusion.^{[ 19 ]} Knowing this critical gap, it is essentially important to further explore the origin of such inhomogeneous kinetic deterioration that strongly correlates with the H2‐H3 phase transition in the cycled Ni83 cathode during high voltage charge.

2.3. Structural‐Heterogeneity Impeded Mass‐Transfer Kinetics

Intergranular cracking and surface reconstructions are widely regarded as the causes of decreased overall diffusion. However, our observations and diagnosis suggest a low likelihood of cracking‐driven diffusion limit. First, large intergranular cracking may result in the isolation of active material, and mass transfer is limited exclusively when cracks are enclosed at the internal or unexposed to the electrolyte. Yet, it is observed that some intergranular cracks propagate throughout the entire secondary particle at high potentials (Figure 2g), leading to electrolyte infiltration with reduced diffusion length. Second, our post‐GITT results (Figure 3a) present the non‐synchronized behaviors of inhomogeneous reduction in the overall “diffusion” kinetics and nearly homogeneous electrical resistance during the charging and discharging processes. There is also a large discrepancy between the dramatic reduction in overall diffusion (5.8 × 10⁻³ times lower at 4.6 V) and a relatively mild electrical impedance hike (4.1 times higher at 4.6 V). These suggest the inhomogeneous reduction in mass‐transfer kinetics is irrelevant to the isolation of active material or cracking. Third, the kinetically impeded H2‐H3 phase transition over cycling (Figure 4f) results in less lattice volume change to drive intergranular cracking. In other words, the intergranular cracking is also kinetically passivated over cycling, as no observation of severe particle pulverization after 1000 cycles at high voltage.

The ever‐growing surface reconstruction over cycling (see STEM images in Figure S15 (Supporting Information) and XAS Ni L‐edge spectra in Figure S16, Supporting Information) is widely regarded as the root of diffusion kinetic deterioration, due to its ionic‐isolating characteristics. However, it is found that the mass‐transfer kinetic deterioration mechanism is not exclusively ascribed to the chemical diffusion reduction. For instance, in the GITT analysis, the chemical diffusion behavior should exhibit the linear relationship between voltage gain and the square root of pulse time (t^0.5 ) in the galvanostatic current pulse, following Fick's 2nd law of diffusion.^{[ 20 ]} The pristine Ni83 cathode in 1st cycle exhibits a neat linear relationship at both middle SOC of 3.8 V and high SOC of 4.3 V (see Figure 5a,b), suggesting that the chemical diffusion principles are well adopted in this material.^{[ 20b ]} In comparison, the Ni83 cathode after 1000 cycles at 4.6 V presents a distinct behavior with initially non‐linear voltage gain (Et^* , i.e., the difference between total voltage gain and the linear voltage gain) and subsequentially linear voltage gain (Et₀ , i.e., the fitted linear slope multiplying with the entire pulse time in t^0.5 ). The chemical diffusion coefficient of the cycled cathode using Et₀ , as shown in Figure 5c, shows approximately homogeneous decay from the pristine cathode in the entire voltage window. A simple analytical diffusion model helps elaborate the statistical reduction of chemical diffusion in bulk cathode material as a function of the thickness of the ionic‐isolating rock‐salt surface layer (Figure 5d). A good agreement of the same‐magnitude reduction in the chemical diffusion, ≈1 × 10⁻⁹ cm⁴ S⁻¹, is shown between the analytical solution for more than 30‐nm thick rock‐salt layer observed via STEM (Figure S15, Supporting Information) and the empirical GITT measurement using the linear Et₀ (Figure 5c) for cycled Ni83 cathode at 4.6 V after 1000 cycles. It is suggested that the ionic‐isolating characteristics of the rock‐salt surface layer led to the homogeneous chemical diffusion reduction over the entire SOC range in the cycled cathode material. Therefore, the inhomogeneous reduction in the overall mass‐transfer kinetics at high SOC, like 1 × 10⁻¹¹ cm⁴ s⁻¹ in the cathode at 4.6 V, is associated with the initially non‐linear Et^* voltage gain as a kinetic potential barrier before chemical diffusion behavior (Figure 5c).

Figure 5.

Severe Kinetic Inhibition at High SOCs in Affiliation to Structural Heterogeneity with Transformed Rock‐Salt Structure in Surface Layer. The comparison of voltage gains in the GITT pulse steps at a) low SOC of 3.8 V and b) high SOC of 4.3 V as a function of the square root of time in the pristine Ni83 cathode and the cycled Ni83 cathode at 4.6 V for 1000 cycles. c) The derived surface‐independent coefficients of chemical diffusion using the linear voltage gain (Et₀) and the overall mass transfer kinetics using the total voltage gain (Et₀+Et ^* ). d) The chemical diffusion coefficient of bulk cathode material, including the rock‐salt surface layer and the “core” layered structure, is analytically calculated as a function of the rock‐salt layer thickness using the experimentally determined diffusion coefficients (4 × 10⁻¹¹ cm² s⁻¹ for NiO structure^{[ 24 ]} vs 3.9 × 10⁻⁸ cm² s⁻¹ for layered structure from GITT test of pristine cathode), respectively. The inlet schematically shows this analytical model of Li diffusion in cathode primary particle of 100 nm radius with a CEI layer (including rock‐salt phase) at the surface. e) The diagrams of layered structure transition in the delithiation process and the rock‐salt structure illustrate two major structural discrepancies of large mismatch strain perpendicular to layer planes and the layer plane shift due to changing stacking order in H2‐H3 transition at high SOC region. f) Synchrotron‐based in‐situ XRD of pristine Ni83 cathode in 1^st cycle. g) The evolution of lattice constant c and a via Rietveld refinements. h) The correlation plots of voltage‐dependent mismatch strain of rock‐salt and layered structures, the “equilibrium” phase transition in GITT‐dQdV, and the non‐linear voltage gain (Et ^* ) standing for the energy barrier of mass‐transfer kinetics.

It should be first noted that the pristine cathode with a clean surface possesses almost homogeneous kinetics between H2‐H3 transition and other transitions at low SOC (Figure 5c), indicating free accommodation of substantial lattice shrinkage in the c‐axis and layer plane gliding for different stacking orders at high SOC. Thus, it is reasonably speculated that the inhomogeneous kinetic reduction strongly correlates to the electrochemical disturbance by undesired structural degradation in the Ni83 cathode during operation at high voltages. Here, we pay special attention to the structural discrepancy between the electrochemically inactive rock‐salt structure and the delithiated phases in a layered structure. Ni‐rich cathodes typically have a hexagonal lattice (denoted as H1) where transition metal ions and lithium ions alternatively occupy within the oxygen framework of octahedral (O3) coordination with the AB CA BC stacking order (Figure 5e). As Li⁺ initially departs from the lattice, the cathode becomes a possible monoclinic phase (M) at ≈20% SOC due to the Jahn‐Teller effect,^{[ 21 ]} and the oxygen framework remains O3 stacking order with enlarged layer spacing (as the increased lattice constant c via in situ HEXRD in Figure 5f). With more Li vacancy, the cathode returns to a hexagonal lattice (H2) at ≈60% SOC with slight lattice shrinkage. Above 75% SOC, there is an abrupt lattice shrinkage in the c‐axis of the hexagonal lattice (H3), like‐3.36% at 4.6 V in Figure 5g, and the structure changes to O1 coordination with a different stacking order of AB AB AB. After cycling at high voltage, the cathode material bears exacerbated rock‐salt‐like surface reconstruction with obvious structural heterogeneities, as illustrated in Figure 5e. Given the fact that the layer spacing along the c‐axis or (003) direction in layered structure is generally smaller than that of inactive rock‐salt structure in (111) plane (d_layer = 2 × d_F(111) = 4.824 Å), their lattice mismatch tends to generate a positive strain at their interface (Figure 5h). Even worse, the abrupt lattice shrinkage at the H2‐H3 transition suffers substantial mismatch strain, like 3.9% at 4.6 V vs 0.26% at 4.0 V. The lattice mismatch represents a stress‐induced energy barrier in inhibiting phase transitions in the Ni83 cathode, which also highly synchronizes with the dramatic increases in the kinetic potential barrier (Et^* in post‐GITT measurements) in Figure 5h. The ever‐growth of mechanically rigid rock‐salt surface layer causes larger Et^* , like 0.011 V at 100 cycles vs 0.051 V at 1000 cycles at 4.6 V, to kinetically impede phase transitions in a delithiated layered structure. Meanwhile, the change of stacking orders in the transition to the H3 phase (i.e., from AB CA BC to AB AB AB ordering) is accompanied by certain gliding between layer planes but is also mechanically pinned by the inactive rock‐salt structure with AB CA BC stacking order in the (111) direction in Figure 5e.

Worthfully noting, this strain‐induced potential barrier (Et^* ) also substantially delays the time response for chemical diffusion, showing up to 400 s in the Ni83 cathodes after 1000 cycles at 4.6 V (see the pulse steps in post‐GITT tests in Figure S17, Supporting Information). This corresponds to a new electrochemical impedance at a very low frequency of <0.004 Hz after cycling (see electrochemical impedance spectroscopy of pristine and cycled cathodes in Figure S18, Supporting Information). Wide variation in time or frequency largely depends on the SOC. This indicates that the cycled cathode materials need more response time to overcome a larger energy barrier from the mismatch strain. The physical frequency characteristics of kinetic energy barrier Et ^* should belong to the range of mass‐transfer kinetics (below 0.5Hz) rather than that of charge‐transfer kinetics in NMC cathodes (100–0.1 Hz).^{[ 22 ]} In addition, the increased charge‐transfer impedance after cycling resulted from poor electronic conductivity of the surface rock‐salt layer by hindering electron transfer (Figure S18, Supporting Information) but is much less significant than the strain‐induced barrier at lower frequency for severe kinetic deterioration at high SOC. The nearly identical frequency of charge‐transfer impedance after cycling reveals the sustained faradaic reaction rate in the cycled cathode, as supported by the highly thermodynamical activity and most layered structure in the cycled cathode after high‐voltage operation. Therefore, the structural heterogeneity with the inactive rock‐salt surface layer, including the mismatched layer spacing and stacking order, could lead to a large mass‐transfer kinetic barrier to inhibit H2‐H3 transition during cycling at high SOC. It has been similarly reported in the condensed matter field that, compared with a single and homogeneous lattice, the heterogeneous lattices or disordering substantially affect diffusion or mass‐transfer behaviors.^{[ 23 ]} Elucidated by the kinetics perspective, severe performance deterioration at high voltage is predominantly ascribed to mechanical‐kinetic inhibition by undesired surface reconstructions, rather than the bulk structural degradation.

As of great interest, current wisdom about formation mechanisms and mitigation strategies of surface reconstruction is briefly discussed to provide a deeper understanding of boosting cathode development. Surface reconstruction in Ni‐rich cathode material mainly underlies undesired phase transformation from layered structure to rock‐salt structure. This transformation is physically rooted in both Ni reduction and lattice oxygen vacancy at the surface of the cathode material. Oxygen vacancies in the close‐packed structure provide a low energy pathway to trigger the migration of Ni²⁺ cations in octahedral sites to the neighboring Li⁺ tetrahedral sites for their similar ionic radii,^{[ 25 ]} eventually forming rock‐salt phase (NiO). The evolution of lattice oxygen vacancy involves complicated mechanisms. On the one hand, the lattice oxygen is thermodynamically prone to release at high SOCs, due to the reduced formation energy of oxygen vacancy from 1.8 eV at the full lithiation to 0.35 eV at the deep delithiation by DFT simulations.^{[ 26 ]} Yet, a large kinetic barrier (2.4 eV) for oxygen migration in the bulk cathode materials substantially limits the oxygen evolution at surface regions.^{[ 26 ]} On the other hand, as many observations of rock‐salt transformation in electrolyte exposure exclusively^{[ 2d ]} or modulated by various electrolytes,^{[ 18} , ^{27 ]} it has been recently unveiled that surface rock‐salt transformation has a strong affiliation to parasitic reactions at the cathode/electrolyte interface.^{[ 28 ]} Parasitic reactions, stemming from electrolyte oxidation at the cathode/electrolyte interface, create a chemically acidic environment to etch away lattice oxygen and offer electrical charge to reduce the Ni valence state.^{[ 28a ]} The rate of oxidative parasitic reactions is nearly exponentially increasing as elevating potentials via high‐precision leakage current (HpLC) measurements (Figure S19, Supporting Information), indicate the aggravated chemical modulation of surface reconstructions at higher voltages.

Pioneering mitigation strategies have been developed accordingly to enhance interface stability and sustain high performance. For instance, element doping at the cathode surface has been effectively adopted to improve the thermodynamic stability of the cathode structure, with the main merit of enhancing the bond strength with lattice oxygen to mitigate oxygen vacancy.^{[ 29 ]} Regulating chemical reactivity at the cathode/electrolyte interface is another key philosophy to mitigate surface reconstructions, and various strategies have been demonstrated along the chemical reaction pathways.^{[ 30 ]} Novel electrolyte or electrolyte additives can either suppress the generation of acidic proton‐bearing species to prevent lattice oxygen loss or possess higher highest‐occupied molecular orbital to improve interface stability.^{[ 31 ]} Artificial interface layers via surface modifications can physically prohibit direct contact with the electrolyte and inhibit the lattice oxygen release from the cathode.^{[ 32 ]} One should note that, in practice, the state‐of‐the‐art doping and coating strategies introduce the electrochemical inactive material into/onto the cathode material, usually leading to certain initial capacity loss. From the perspectives of structure‐heterogeneity‐induced kinetic inhibition in this study, it is suggested to revisit whether the introduction of electrochemical inactive material is kinetically unfriendly to accommodate the lattice change during phase transitions. Likewise, more attention should be paid to the performance disturbance from the possibly distinct kinetic characteristics between heterogeneous structure domains in other cathode materials. Herein, we believe that the kinetic perspective highlighted in this study might offer more insights into the state‐of‐art mitigation strategies for boosting technology innovation.

3. Conclusion

We have demonstrated the superior importance of kinetic characteristics in Ni‐rich cathode during long‐term cycling at high voltage. Through the deconvolution of irreversible and reversible capacity loss, it is found that the predominant contribution from kinetic limits unpuzzles the contradiction of severe performance deterioration and insignificant bulk structural degradation. The so‐called diminished H2‐H3 phase transition in the responsibility of severe capacity decay in high‐voltage Ni‐rich cathode practically exhibits high reversibility under free kinetic circumstances, which run counters to the prevailing beliefs on irreversible bulk fatigue and lattice collapse degradation. This well elaborates that the degraded H2‐H3 transition is the limited observation in regular galvanostatic or constant‐current tests that presume the homogeneous kinetics among all phase transitions during cycling.

This study also finds that the inhomogeneous kinetics reduction in the cycled Ni‐rich layered cathode at high voltage cannot be explained by the prevailing understanding, including chemical diffusion limit by surface reconstruction and contact loss by cracking. By contrast, it exhibits a strong correlation with the mismatched lattice strain between the delithiated layered structure and the inactive rock‐salt phase. The perspectives of kinetic loss due to structural heterogeneity can be applied to revisit the well‐known initial capacity loss by coating or doping mitigation strategy of introducing inactive and rigid materials and also investigate challenges in other cathode chemistries with heterogeneous structure domains. This study reveals that kinetics is fundamentally vital to bridge micro‐structural heterogeneity and bulk phase transition, and also practically important to offer key insights toward long‐term cycling and potentially fast‐charging challenges to boost future technology innovation for unleashing high energy‐density Ni‐rich layered cathodes.

4. Experimental Section

Material Synthesis

The precursor for making the LiNi_0.83Mn_0.1Co_0.01O₂ (Ni83) cathode had a nominal composition of Ni_0.83Mn_0.1Co_0.07(OH)₂ and was synthesized via a co‐precipitation method. To obtain the layered Ni83 powders, the Ni_0.83Mn_0.1Co_0.07(OH)₂ was first mixed with LiOH·H₂O by acoustic mixing. The mixture was then preheated at 500 °C for 12 h and then calcined at 790 °C for another 12 h in O₂. During the heating process, the temperature was increased at a constant rate of 2 °C min⁻¹. To compensate for the lithium loss during high‐temperature heat treatment, an extra 5% in excess of the stoichiometric requirement of lithium source (LiOH·H₂O) was added.

Electrode Preparation

Cathode laminates on Al foils were prepared by mixing 92 wt.% Ni83 powder with 4 wt.% carbon black conductive agent and 4 wt.% polyvinydene fluoride (PVDF) binder in 1‐methyl‐2‐pyrrolidone (NMP) solvent. The well‐mixed slurry was then cast on Al foil using a 225 µm doctor blade. The cathode laminates were dried at 75 °C in air overnight and then calendared to ≈33% porosity with electrode material thickness of 34 µm thick. The mass loading of the cathode (9.7–10.7 mg cm⁻²) was strictly selected for different voltage windows to match with the reversible capacity graphite anode to achieve the target N:P ratio of 1.10 (please see Figure S2 (Supporting Information) with the detailed preparation in the Supporting Information). The graphite anodes were made in a dry room at the NorthEast Center for Chemical Energy Storage (NECCES) at Binghamton University with less than 0.5% relative humidity. Anode laminates were prepared by mixing 96 wt.% synthetic graphite, 1 wt.% carbon black, 1 wt.% carboxymethyl cellulose (CMC), and 2 wt.% styrene‐butadiene rubber (SBR) in deionized water using a Media Tech Co. planetary dispersing mixer. The resulting slurry was coated onto copper foil using a Media Tech Co. coating machine. The target mass loading of the graphite anode was ≈6.8 mg cm⁻² for the area‐specific capacity of 2.31 mAh cm⁻², followed by a calendering process to 42.5 µm thickness. All laminates were punched into discs (14 mm in diameter for cathodes and 15 mm for anodes), and the punched cathodes and anodes were vacuum dried at 110 and 130 °C overnight, respectively, before cell assembly. CR2032 coin cell parts were used to assemble full cells (NMC and graphite) and half cells (NMC and 50 µm Li metal counter), with a Celgard 2320 separator. The liquid electrolyte was 40 µL of 1.2 m LiPF₆ in ethylene‐carbonate (EC)/ethyl‐methyl‐carbonate (EMC) (3:7 by weight) from Tomiyama Pure Chemical Industries, Ltd.

Cell Assembly and Electrochemical Tests

Electrochemical cycling evaluation of full cells using Ni83 cathode and graphite anode was conducted on a Landt battery testing system (CT3001A‐5V10mA) at room temperature. To better elucidate the physical root of performance loss on the cathode, it is primarily necessary to differentiate the influence of the cathode and the anode. Generally, graphite anode holds a relatively steady working potential at ≈0.05 V, while the Ni‐rich cathode possesses multiple potential plateaus and usually exhibits poor behavior at a high potential of > 4.3V. In this study, the upper potentials of the Ni83 cathode were determined at 4.2, 4.4, and 4.6 V for failure mechanism investigation. Restricting the N:P ratio of 1.10 (i.e., the ratio of practical capacities between the graphite anode lithiating to 0.05 V and the cathode delithiating to targeted potential) approximately assures the same working potential of graphite anode and utilizing the various SOCs of the cathode in full cells. Hereby, the upper voltages of 4.15, 4.35, and 4.55 V in full cells roughly correspond to the cathode's upper potential at 4.2, 4.4, and 4.6 V in half cells, to ensure the same delithiation state of cathodes in the full cells and subsequentially the harvested half cells. The cycling test protocol includes two formation cycles at C/10 current rate, then 1000 cycles at 1C current rate (1C = 214 mA g⁻¹) with the constant‐current constant‐voltage (CCCV) mode in the charging (constant voltage for 1 h) and constant current (CC) mode in the discharging, followed by 10 cycles at C/10. Two different voltage windows, 2.5–4.15 V and 2.5–4.35 V, were selected to evaluate the electrochemical performance of Ni‐rich NMC/graphite full cells. Electrochemical impedance tests have been carried out on the BioLogic SP‐300 system in the frequency range of 7–3 MHz with a voltage magnitude of 10 mV.

A home‐built high‐precision leakage current (HpLC) system based on Keithley 2401 source meters was used to measure the steady leakage current as a function of charging potential.^{[ 33 ]} The steady leakage current at each potential was obtained by fitting the decaying current profile with an exponential decay function (y = A × exp(‐x/t) + y₀).^{[ 32} , ^{33 ]} Galvanostatic intermittent titration technique (GITT) measurements were performed in half cells using Li metal anode. Before GITT tests, half cells first followed two formation cycles at C/10 in the window voltages and then discharged to 2.5 V at C/50 to permit the complete lithiation into cathode materials. Subsequentially, galvanostatic current pulses of C/50 rate (1C = 200 mA g⁻¹) were applied for 10 min before a 50‐min relaxation step. The charging and discharging processes terminated when the potential at the end of the relaxation step exceeded the upper limit or the lower limit of voltage windows, respectively. The diffusion coefficient of Li⁺ (D_Li ⁺) can be obtained in the following equation:

$$D_{L{i}^{+}}=\frac{4}{\pi \tau }{\left(\frac{m_B{V}_M}{M_{B}S}\right)}^2{\left(\frac{\mathrm{\Delta }{E}_s}{\mathrm{\Delta }{E}_t}\right)}^2$$ (1)

where τ is the time of galvanostatic current pulses, ΔE_t is the change of the transient potential before and after the current pulses, ΔE_s is the change of the steady state voltage, and S is the contact area between the active material and electrolyte. V_M , M_B , and m_B are the molar volume, molar mass, and mass of the cathode material, respectively. The analytical chemical diffusion model of a primary cathode particle with 100 nm radius in Figure 5d was calculated based on the total diffusion length (L_layered + L_rock‐salt = 100 nm) divided by the sum of individual diffusion time (t_total), i.e., the “core” layered structure using the GITT‐measured diffusion coefficient (3.9 × 10⁻⁸ cm² s⁻¹) in the pristine cathode and the rock‐salt surface layer experimentally determined diffusion coefficient (4 × 10⁻¹¹ cm² s⁻¹ for NiO structure^{[ 24 ]}), as the following equation:

$$D_{bulk}\approx \frac{{\left(L_{layered}+L_{rock-salt}\right)}^2}{t_{total}}$$ (2)

$$t_{total}=\frac{{\left(L_{layered}\right)}^2}{D_{layered}}+\frac{{\left(L_{rock-salt}\right)}^2}{D_{rock-salt}}$$ (3)

Material Characterizations

Microscopic images of electrode materials and laminates were obtained from a Hitachi S‐4799‐II scanning electron microscope (SEM) with a Brucker energy dispersive X‐ray spectroscopy (EDX), a FEI Talos F200X scanning transmission electron microscope (STEM) with EDX, STEM, and a JEOL JEM2100F high‐resolution transmission electron microscope (HRTEM) with selected area electron diffraction (SAED) at the Center for Nanoscale Materials (CNM) at Argonne National Laboratory. Cross‐sectioned NMC secondary particles were etched by a focused ion beam (FIB) integrated with the SEM, with a thinning procedure at 30 kV and 40 pA. Synchrotron high‐energy X‐ray diffraction (HEXRD) was carried out for NMC cathodes at beamlines 11‐ID‐C (0.1173 Å wavelength) and 17‐BM (0.45185 Å) of the Advanced Photon Source (APS) at Argonne National Laboratory. Synchrotron X‐ray fluorescence (XRF) on graphite anodes and hard X‐ray absorption spectroscopy (hard XAS, including near‐edge and extended spectrum of transition metal K‐edge) on NMC cathodes were performed at beamlines 8‐BM‐B and 20‐BM‐B of the APS, respectively. Synchrotron XRF permits the large‐area scanning (2 × 2 cm²) in the transmission mode to quantify the total amount of migrated transition metals in the entire graphite anode. The measurements of Ni K‐edge, Co K‐edge, and Mn K‐edge were performed at room temperature in transmission mode. For the XRF experiments, samples were raster‐scanned with a focused 10 keV X‐ray beam in a 25 µm step size. The 2D projection data were then processed and quantified by XRF‐MAPS.^{[ 34 ]} Soft XAS for Ni L‐edge on cycled cathodes were performed at the Advanced Light Source (ALS) at the Lawrence Berkeley National Laboratory (LBNL), including total electron yield (TEY) and total fluorescence yield (TFY) channels. Soft XAS collects surface and relatively bulk material information simultaneously through the total electron yield (TEY) channel with a few‐nanometer electron escape depth and the total fluorescence yield (TFY) channel with a hundred‐nanometer photon penetration depth, respectively.^{[ 35 ]} The Ni²⁺ contents in cycled cathode samples were fitted with the linear superposition of Ni spectra at different valence states. High‐pressure ion chromatography (HPIC) was carried out on a Dionex ICS‐6000 system with an UV detector to determine the amount of transition metal contents in the collected electrolyte solution. After dissembling in the Ar‐filled glovebox, the cell components are thoroughly washed with 1 mL of acetonitrile (HPLC grade) and then diluted with 1 mL ultra‐pure water to obtain the solution. The IC system was calibrated with calibration standards (ultra‐pure water dissolving trace Mn, Co, or Ni ions).

Conflict of Interest

The authors declare no conflict of interest.

Supporting information

Supporting Information

Cai J., Zhou X., Li L., Yang Z., Huang X., Li J., Wang G., Zhu Q., Li T., Sun C.‐J., Zhuo Z., Suzana A., Bai J., Gudavalli G., Karami N., Chernova N. A., Upreti S., Prevel B., Yang W., Liu Y., Xu W., Chen Y., Song S., Zhang X., Wang L., He X., Wang F., Xu G.‐L., Chen Z., Kinetically Dormant Ni‐Rich Layered Cathode During High‐Voltage Operation. Adv. Mater. 2025, 37, 2419253. 10.1002/adma.202419253

Data Availability Statement

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