Flash upcycling of spent LiCoO_2 into oxygen-suppressed lithium-replenishing agent for high-performance batteries

Domain: physics
Published: 2026-01-05
Editorial outcome: minor_revision

Abstract
Battery recycling is essential to mitigate resource depletion and improve environmental sustainability. However, conventional metallurgical and direct regeneration methods involve trade-offs among energy input, environmental impact, and feedstock adaptability. Here we report a universal upcycling method that converts spent LiCoO2 into Li6CoO4 within 10 seconds via flash Joule heating, achieving complete lithium and cobalt recovery regardless of degradation state. The resulting Li6CoO4 serves as a high-capacity sacrificial additive to offset active lithium loss. To overcome its inherent air sensitivity and oxygen release during delithiation, a conformal sulfur coating was applied to stabilize surface chemistry and redirect oxygen evolution into sulfate formation. This sulfur-mediated mechanism effectively suppresses gas generation and parasitic side reactions, enabling non-destructive lithium replenishment. When integrated into graphite||LiFePO4 pouch cells, the stabilized Li6CoO4 achieves 91.4% capacity retention over 1400 cycles at a current density of 80 mA g−1. Comparative life-cycle and techno-economic analyses reveal clear reductions in energy consumption and CO2 emissions, along with improved economic returns. This work provides a scalable route that bridges positive electrode waste recovery with high-performance lithium supply, advancing closed-loop battery systems for sustainable energy storage.

Introduction
LiCoO2 (LCO) is the primary positive electrode material used in rechargeable batteries for portable consumer electronics1. With the widespread use and frequent replacement of these devices, over 100,000 tons of spent LiCoO2 (SLCO) batteries are discarded annually worldwide, highlighting the urgent need for sustainable waste management solutions2,3. Current recycling strategies for lithium-ion batteries (LIBs) can be broadly categorized into metallurgy-based elemental extraction and direct regeneration4,5. The former includes pyrometallurgy, hydrometallurgy, and their combinations, in which the chemical bonds of electrode materials are broken under high-temperature or acidic/alkaline conditions to extract the target metal elements6. These methods offer high tolerance for feedstock heterogeneity but compromise economic and environmental sustainability, and are therefore typically regarded as “downcycling” strategies. Additionally, valuable lithium is often lost in the mixed slag or through evaporation at high temperatures, further limiting their overall efficiency7.

To address these drawbacks, direct regeneration has emerged as a more sustainable alternative, aiming to restore the functional properties of spent electrodes with minimal structural disruption8,9. Typical direct regeneration methods employ molten salt methods, hydrothermal treatments, or solid-state sintering to reconstruct the crystal lattice and revive electrochemical activity10. This method is particularly well-suited for materials with stable crystal structures, such as olivine LiFePO4 (LFP), where degradation is primarily due to lithium loss11,12. However, for layered oxide materials like LCO and LiNixCoyMzO2 (NCM), degradation is induced by intertwined chemical, mechanical, and electrochemical interactions, making direct repair considerably more challenging13. While several studies have explored the direct repair of spent LCO into regenerated LCO14,15, achieving stable long-term cycling remains complex. The incomplete performance recovery ultimately results in the inefficient reuse of lithium and cobalt in new cells.

Achieving high resource utilization efficiency requires not only optimized recycling strategies, but also careful consideration of how recovered elements perform during their second life, particularly in long-term cycling stability. In this regard, extending the operational lifespan of LIBs is critical, as it enables fixed lithium inventories to deliver more energy throughout their service life, thereby enhancing overall resource efficiency. Prolonged use not only defers battery retirement but also reduces waste generation and eases the burden of end-of-life battery management. One promising approach to support long-term performance is to supplement cells with an additional lithium supply, which compensates for irreversible lithium loss during cycling and helps minimize capacity fade16–18.

Toward this end, we propose an upcycling strategy that repurposes SLCO into Li6CoO4 (L6CO), a high-value sacrificial lithium-replenishing agent. This approach circumvents the need for complex repair of the layered framework required for stable cycling, but instead reconstructs an antifluorite structure that irreversibly releases active lithium to enhance energy density and extend battery lifespan. L6CO does not function as a cycling positive electrode but becomes electrochemically inactive after initial decomposition19, thus avoiding concerns over long-term stability. The as-prepared L6CO exhibits a low decomposition potential (~3.2 V vs. Li+/Li), a high irreversible capacity (856 mAh g−1), and great rate capability, outperforming commercially available alternatives such as Li2NiO2 (LNO) and Li5FeO4 (LFO) in both capacity and oxidation kinetics20,21. The entire conversion is achieved within 10 s via flash Joule heating (FJH). With the simple addition of a lithium source (Li2O), the residual carbon and polyvinylidene difluoride (PVDF) in the spent electrode are directly utilized as in situ reducing agents during thermal conversion, eliminating the need for external reductants and simplifying the overall process. This method also demonstrates high universality, enabling efficient conversion of SLCO across various degradation states while achieving near-complete recovery of both lithium and cobalt.

Similar to other lithium-rich oxides (e.g., Li2O, Li2O2, and Li5FeO4), L6CO suffers from poor air stability due to its high lithium content22,23. In humid environments, it readily reacts with moisture and CO2 to form LiOH and Li2CO3, leading to slurry gelation and complicating storage and processing. While high-temperature carbon coating is commonly used to enhance air stability, forming a dense carbon layer on L6CO is challenging due to the oxidation of carbon and the reduction of transition metals at elevated temperatures. Moreover, oxygen participates in charge compensation during the delithiation of L6CO. The consequent release of lattice oxygen generates reactive oxygen species and O2 gas24,25, which can trigger parasitic side reactions with the electrolyte and compromise both safety and cycling stability of batteries26,27.

To mitigate these challenges, we introduce a low-temperature sulfur coating strategy to engineer stable L6CO. By leveraging the low melting point and hydrophobicity of elemental sulfur, a uniform and robust coating can be readily formed on L6CO (L6CO@S), without altering the underlying structural integrity. Beyond physical encapsulation, sulfur exhibits high oxygen-trapping efficiency by chemically interacting with lattice oxygen during charging to form stable SO4 2− species28,29, in which a single sulfur atom coordinates with up to four oxygen atoms. This sulfur-mediated anchoring mechanism effectively suppresses O2 evolution and mitigates parasitic side reactions. In contrast to conventional carbon coatings that merely act as physical barriers, sulfur provides a chemically active and more effective pathway to stabilize L6CO during initial delithiation.

The as-synthesized L6CO@S composite functions as an effective lithium compensator, offering improved air stability, enhanced processability, and suppressed oxygen evolution. After Li supply with L6CO@S, graphite||LiFePO4 (Gr||LFP) full cell demonstrates enhanced cycling stability, achieving 91.4% capacity retention after 1400 cycles at 0.5 C (1 C = 160 mA g−1), compared to 72% retention in control cells without L6CO@S. It also enhances the capacity retention of graphite||LiCoO2 (Gr||LCO) cells by 20% over 300 cycles, demonstrating broad applicability across different battery systems. Moreover, utilizing SLCO as a low-cost precursor helps mitigate cobalt-related cost concerns and supports the economic viability of this strategy. Life cycle analysis confirms the reduction in energy consumption and CO2 emissions, along with improved economic profitability, underscoring the practicality of the upcycling strategy. Together, these findings establish a sustainable route for the high-value reuse of spent batteries in next-generation LIBs.

Results
Rapid upcycling of SLCO into L6CO
The SLCO electrodes were retrieved from degraded 2–10 Ah LCO pouch cells (Supplementary Fig. 1). The electrode materials were efficiently separated from the aluminum current collectors using the water electrolysis-induced separation (WES) method we previously reported30. The collected SLCO pieces were then ground into a fine powder before further processing. The inductively coupled plasma-optical emission spectrometry (ICP–OES) analysis further confirmed that the residual Al content in the SLCO powder was only 1.8 mg L−1, and no Cu impurity was detected (Supplementary Fig 2). Figure 1a and Supplementary Fig. 3 present the scanning electron microscopy (SEM) morphology of the SLCO material, revealing severe interlayer slippage and microcrack formation on the particle surfaces. These structural degradations are primarily caused by prolonged lithium depletion and mechanical fatigue during repeated Li-ion intercalation/deintercalation cycles. High-resolution transmission electron microscopy (HRTEM) analysis (Fig. 1b) shows that while the bulk SLCO maintains its layered structure, the surface regions undergo partial phase transformation into a disordered spinel structure, with lattice fringes corresponding well to Co3O4. X-ray diffraction (XRD) analysis of SLCO (Fig. 1c) further confirms this structural evolution, revealing distinct diffraction peaks corresponding to both the layered LCO phase (PDF# 75-0532) and the spinel Co3O4 phase (PDF# 43-1003). Quantitative phase analysis via Rietveld refinement (Supplementary Table 1) indicates that the spinel component accounts for approximately 10 wt.%. The formation of densely distributed Co3O4 domains at the particle surface markedly impedes lithium-ion diffusion and increase the thermodynamic barriers for the direct restoration of SLCO into new LCO.

Elemental analysis by ICP–OES quantifies a 12% lithium deficiency in SLCO (Li/Co=0.882) relative to stoichiometric LiCoO2. The degraded material exhibits severely compromised electrochemical performance in half-cells, with an initial capacity of only 67.3 mAh g−1 and rapid capacity fading during cycling (Supplementary Fig. 4). These observations indicate that capacity loss is not solely attributable to lithium depletion but also reflects extensive structural degradation. While high-temperature sintering with stoichiometric lithium compensation partially restores crystallinity and capacity, the regenerated material remains inferior to commercial LCO (CLCO), particularly in terms of cycling stability (Supplementary Figs. 5 and 6). These findings underscore the limitations of conventional direct regeneration approaches when applied to severely degraded SLCO, where irreversible phase transformations and structural collapse prevail as dominant degradation mechanisms.

The direct conversion of SLCO to L6CO eliminates the need for complex reconstruction of a precise layered structure. In this approach, SLCO was uniformly mixed with stoichiometric Li2O and subjected to vacuum sintering via a one-step FJH process. Systematic optimization across a temperature range of 800–1300 °C revealed a clear trend in phase evolution. XRD analysis (Supplementary Fig. 7) shows the gradual disappearance of SLCO-related peaks and the concurrent intensification of L6CO peaks with increasing temperature, indicating improved reaction completeness. At 1200 °C, a 10 s pulse enables complete conversion into phase-pure Li6CoO4, whereas further increases in temperature or dwell time do not improve the transformation and instead incur unnecessary energy input. Thus, 1200 °C for 10 s was identified as the optimal synthesis condition, balancing transformation kinetics with energy efficiency. The corresponding temperature–time profile during the FJH process (Supplementary Fig. 9) shows that the sample temperature rises sharply to 1200 °C, confirming the fast heating characteristics of this technique.

The resulting product exhibits a distinct blue coloration (Supplementary Fig. 10). HRTEM analysis identifies well-ordered lattice fringes with an interplanar spacing of 0.379 nm, corresponding to the (101) planes in the antifluorite structure (Fig. 1d and Supplementary Fig. 11). XRD analysis and Rietveld refinement (Fig. 1e and Supplementary Table 2) reveal a phase-pure L6CO diffraction pattern, with complete elimination of LCO and Co3O4 signals and no detectable impurities, demonstrating the effectiveness of the FJH approach for rapid and complete phase conversion.

Our FJH upcycling strategy offers a distinct advantage over conventional direct recycling methods by eliminating the need for complete removal of PVDF binders and conductive carbon prior to regeneration31. Instead, these residual components are strategically utilized, serving as in situ reducing agents that facilitate efficient cobalt reduction during the formation of L6CO. This eliminates the need for both extensive pre-treatment steps and external reductants. Microstructural characterization (Supplementary Fig. 12) confirms that the high temperatures during FJH effectively decompose residual carbon and PVDF in pristine SLCO. The resulting L6CO exhibits a smooth single-crystal morphology, indicating the complete removal of surface contaminants. FTIR of SLCO displays characteristic absorption peaks corresponding to PVDF (900–1400 cm−1), which almost completely disappear after FJH (Supplementary Fig. 13), confirming its full decomposition. Owing to the brief reaction duration, the released fluorinated volatiles are transient and concentrated in the off-gas stream, which makes them amenable to engineering control. This self-contained upcycling protocol clearly simplifies battery recycling workflows while maintaining stringent purity standards. EDS elemental mapping shows a homogeneous distribution of cobalt and oxygen throughout the particles, confirming compositional uniformity (Supplementary Fig. 14).

Electrochemical evaluation demonstrates promising lithium compensation capability of L6CO (Fig. 1f), with an initial charge capacity of 856.2 mAh g−1 at 0.1 C (1 C = 800 mA g−1). The voltage profile exhibits two well-defined oxidation plateaus at 3.25 V and 3.65 V (vs. Li+/Li), with negligible discharge capacity (<2.5% of charge capacity), characteristic of an efficient lithium reservoir. Cyclic voltammetry (CV) profiles further corroborate these electrochemical characteristics, showing great alignment between oxidation peaks and galvanostatic plateaus (Supplementary Fig. 15). Notably, L6CO maintains high rate performance (Supplementary Fig. 16), delivering 759 mAh g−1 at 0.5 C (88.6% retention compared to 0.1 C), outperforming commercial LFO references, which retain only 57.6% of their capacity under identical testing conditions (Supplementary Fig. 17). The enhanced oxidation kinetics enable complete decomposition during the first cycle and suppress parasitic reactions during subsequent cycling.

Unlike conventional regeneration methods that struggle with the variable degradation of different SLCO materials, our approach demonstrates great versatility. We systematically evaluated three SLCO samples with varying degradation levels from different spent batteries, with initial capacities ranging from 20.5 to 100.3 mAh g−1 and corresponding Li/Co ratios between 0.774 and 0.882. Remarkably, regardless of the initial degradation state, all samples were successfully transformed into phase-pure L6CO with precise stoichiometric control, as confirmed by ICP–OES analysis (Supplementary Fig. 18). Additionally, the upcycled products exhibited consistent electrochemical performance, each delivering charge capacity of approximately 850 mAh g−1 (Fig. 1g). To simulate a broader range of degradation conditions, fresh LCO was blended with varying amounts of Co3O4 (5–50 wt.%) and a controlled amount of carbon black was added as the reducer. Despite these deliberate compositional variations, FJH consistently yielded phase-pure L6CO (Supplementary Fig. 19). This consistency underscores the robustness of our method for practical application across diverse SLCO feedstocks.

Figure 1h schematically illustrates our closed-loop upcycling strategy, where the regenerated L6CO acts as a high-efficiency Li supply additive. When combined with fresh positive electrode materials, it enables lithium inventory exceeding 100% in next-generation battery systems. This approach simultaneously extends battery service life and maximizes energy output from limited lithium resources, demonstrating high resource efficiency for both upcycling and reuse.

In contrast to conventional regeneration methods, which require precise phase control and customized calcination protocols for different feedstocks but still fail to achieve complete electrochemical recovery and high resource recovery efficiency, our method offers clear improvements. Quantitative comparisons (Fig. 1i and Supplementary Table 3) demonstrate that our approach outperforms conventional recycling in both material recovery and temporal efficiency. The rapid thermal kinetics of FJH effectively suppress lithium volatilization and enabling near-100% recovery of both lithium and cobalt (Supplementary Fig. 20), while reducing processing times by over an order of magnitude. Together, these advantages establish a grave–to–cradle pathway that unites sustainable upcycling with practical reusability.

Engineer stable L6CO through sulfur coating
The high lithium content of L6CO renders it inherently unstable in ambient air. Its highly alkaline surface readily reacts with moisture, leading to gelation of electrode slurries during electrode fabrication. Upon oxidation, each mole of L6CO decomposes to release five moles of Li+ and one mole of O2, forming a solid disordered-LiCoO2 residue, as demonstrated by Eq. (1):

Li6CoO4 → 5Li+ + LiCoO2 + O2 (1)

The substantial O2 evolution poses safety challenges and induces undesirable side reactions. To address these drawbacks and enhance practical processability, a sulfur coating strategy was employed. Elemental sulfur was uniformly coated onto L6CO via a low-temperature thermal treatment at 160 °C. The resulting L6CO@S exhibits a noticeable color change from blue to green (Supplementary Fig. 21), confirming the formation of a conformal sulfur layer.

SEM images of L6CO@S reveal a smooth and uniform particle surface following sulfur coating, with no observable aggregation of sulfur particles (Supplementary Fig. 22). Transmission electron microscopy (TEM) image coupled with EDS elemental mapping reveals a homogeneous sulfur distribution across the surface (Fig. 2a). Cryo–TEM further confirms a continuous and conformal sulfur coating that fully encapsulates the L6CO particles, with an average thickness of approximately 20 nm (Fig. 2b and Supplementary Fig. 23). XRD patterns (Supplementary Fig. 24) show only the characteristic peaks of crystalline L6CO, with no detectable peaks corresponding to crystalline sulfur, indicating that sulfur remains in an amorphous state while the L6CO core structure is well preserved. Importantly, the sulfur coating does not compromise the electrochemical performance of L6CO. Even with a sulfur content as high as 10 wt.%, L6CO@S still delivers a high initial charge capacity of 627.9 mAh g−1 at 0.5 C (1 C = 600 mA g−1), confirming that lithium extraction remains efficient despite the surface modification (Supplementary Fig. 25). Moreover, galvanostatic intermittent titration technique (GITT) measurements reveal that the Li+ diffusion coefficients (DLi) of pristine L6CO and L6CO@S remain in the same order of magnitude before and after coating (Supplementary Fig. 26), indicating that lithium-ion transport is not markedly hindered by the sulfur layer. In situ XRD analysis during charging up to 4.3 V reveals a progressive attenuation of the (101) and (201) diffraction peaks of L6CO (Fig. 2c), reflecting a transition from long-range crystalline order to a more disordered structure upon lithium extraction.

Improved air stability for L6CO@S. The air stability of both pristine L6CO and L6CO@S was systematically evaluated through gravimetric, structural, and electrochemical analyses after exposure to ambient conditions (relative humidity, 30% RH). The uncoated L6CO also shows a noticeable color transformation from blue to bluish-gray after 72 h of exposure (Fig. 2d), visually indicating chemical degradation due to surface reactions with ambient H2O and CO2. Gravimetric analysis that pristine L6CO undergoes rapid degradation, with a 12% mass gain after 72 h due to the formation of surface hydroxides and carbonates. In contrast, L6CO@S exhibits improved stability, showing only ~1% mass increase under identical conditions (Fig. 2e). XRD patterns of both L6CO and L6CO@S after exposure to 30% relative humidity at different time points are shown in Fig. 2f and Supplementary Fig. 27. After 12 h, L6CO develops prominent peaks corresponding to Li2CO3, indicating severe surface degradation, while L6CO@S shows no detectable degradation products even after 48 h. SEM images in Supplementary Fig. 28 further confirm this contrast: after 24 h in 30% RH air, L6CO develops flaky alkaline surface residues, whereas L6CO@S maintains a clean, smooth particle surface.

Electrochemical testing further underscores the protective effect of the sulfur coating. After 72 h of exposure to 30% RH air, uncoated L6CO retains only 61.2% of its initial capacity, whereas L6CO@S maintains 92.9%, delivering 634.6 mAh g−1 (Fig. 2g and Supplementary Fig. 29). Even under harsher conditions (60% RH for 48 h), L6CO@S retains 88.8% of its original capacity (Supplementary Fig. 30). Beyond performance decay, surface degradation also introduces great processing challenges. After 24 h in 30% RH air, uncoated L6CO induces strong gelation due to its surface alkalinity, while slurry containing L6CO@S remains well-dispersed and flowable (Supplementary Fig. 31). These results highlight the high environmental tolerance and manufacturing compatibility of L6CO@S.

Contact angle measurements demonstrate a clear improvement in water resistance after sulfur coating. The water contact angle increases dramatically from 13.96° for uncoated L6CO to 52.12° for the L6CO@S composite (Fig. 2h), highlighting the coating’s effectiveness in enhancing surface hydrophobicity. This effect is also visually corroborated: water droplets rapidly spread and trigger a blue–to–yellow color change on uncoated L6CO, reflecting immediate surface reactions. In stark contrast, droplets remain nearly spherical on L6CO@S for over 1 minute without inducing any visible color change (Supplementary Fig. 32), demonstrating the coating’s effectiveness in suppressing H2O-driven degradation. We calculated the adsorption energies of H2O and CO2 molecules on the (101) surface of L6CO (Fig. 2i). Both molecules interact through their oxygen atoms coordinating with surface metal sites (Li and Co), resulting in strong adsorption energies of −2.16 eV for H2O and −1.68 eV for CO2. These values highlight the pronounced reactivity of L6CO toward ambient moisture and carbon dioxide. In contrast, when the sulfur coating is applied, the adsorption occurs instead on the sulfur surface layer, where the interactions are considerably weaker. The adsorption energies are reduced to −0.62 eV for H2O and −0.57 eV for CO2, indicating a substantial suppression of surface affinity for these molecules. These striking contrasts highlight the sulfur layer as a robust barrier that shields L6CO from atmospheric exposure.

Suppressed oxygen evolution for L6CO@S. The gas evolution behavior of L6CO and L6CO@S was systematically investigated via online electrochemical mass spectrometry (OEMS). For pristine L6CO, charge compensation during the initial ~30% state of charge (SOC) is primarily achieved through the oxidation of Co2+, with no detectable O2 evolution. Beyond this point, lattice oxygen begins to participate in the oxidation process, resulting in apparent O2 release (Fig. 3a). At higher voltages, a rise in CO2 signals was observed, attributed to electrolyte decomposition triggered by reactive oxygen species such as oxygen radicals (O•), peroxide anions (O2 2−), or singlet oxygen (1O2)32. In contrast, L6CO@S exhibited nearly complete suppression of gas evolution, with O2 and CO2 signals remaining close to background levels and no detectable SO2 formation (Fig. 3b and Supplementary Fig. 33). This complete inhibition of gas evolution underscores the dual protective function of the sulfur layer: stabilizing lattice oxygen and mitigating electrolyte decomposition at high voltages.

The oxidation state of sulfur was further examined using X-ray photoelectron spectroscopy (XPS) analysis of L6CO@S at different SOC (Fig. 3c). Upon charging, the S 2p spectra shift from the characteristic peaks of elemental sulfur (~162.5 eV) to a new peak at ~168.5 eV, corresponding to the S–O bond in SO4 2−. This transformation indicates the involvement of sulfur in charge compensation. XPS depth profiling yielded consistent results across different etching depths (Supplementary Fig. 34), confirming that the transformation is spatially uniform across the particle surface. Importantly, this sulfur oxidation process is not driven by cobalt but is instead closely associated with the presence of reactive lattice oxygen. A control experiment using sulfur-coated LCO, synthesized via the same method, showed negligible sulfur oxidation when charged to 4.3 V (Supplementary Fig. 35). This result is consistent with LCO’s minimal oxygen release at this voltage, further confirming the critical role of lattice oxygen in driving sulfur oxidation. In addition, we evaluated the air stability of the sulfur coating. S 2p XPS spectra of L6CO@S stored in air for 3 and 7 days showed no evidence of sulfate formation (Supplementary Fig. 36), indicating that the sulfur does not undergo spontaneous oxidation in ambient conditions. The oxidation to sulfate in our system requires both electrochemical driving force and the presence of highly reactive oxygen species.

X-ray absorption spectroscopy (XAS) measurements were conducted to gain further insight. The S K-edge X-ray absorption near-edge structure (XANES) spectra reveal the valence evolution of sulfur (Fig. 3d). The pristine L6CO@S electrode exhibits an absorption edge that coincides with elemental sulfur. From 0% to 40% state of charge (SOC), charge compensation is dominated by Co oxidation, and the S K-edge spectrum remains unchanged, indicating no sulfur oxidation. Between ~40% and 60% SOC, a transient feature appears between the positions of elemental sulfur and Li2SO4, consistent with short-lived partially oxidized sulfur species (e.g., sulfites), which disappear at higher SOC. Beyond 60% SOC, the edge shifts toward higher energy, with the emergence of features at 2478.6 eV and 2483.4 eV, characteristic of sulfate species. At full charge, the spectrum nearly overlaps with that of Li2SO4, confirming the complete oxidation of sulfur to SO4 2−. Capillary electrophoresis was further employed to quantify SO4 2− in fully charged L6CO@S electrodes. The measured concentration (~39 ppm) closely matched the theoretical value (Supplementary Fig. 37), providing independent confirmation of complete sulfur oxidation.

The evolution of the Co oxidation state during charge was also examined via Co K-edge XAS. The XANES spectra (Supplementary Fig. 38) show that both L6CO and L6CO@S undergo oxidation from +2 to +4 as delithiation progresses. Notably, at 50% SOC, L6CO@S exhibits a slightly lower Co oxidation state compared to bare L6CO, suggesting partial charge compensation by sulfur. Upon further charging, a shift of the Co K-edge to lower energy is observed, consistent with lattice oxygen oxidation and partial Co–O bond breaking. At 100% SOC (4.3 V), both L6CO and L6CO@S show a Co near-edge shift to lower energy relative to LCO, which can be attributed to further Co–O bond breaking and ligand–to–metal charge transfer, leading to the formation of Co3O4-like phases with a lower average Co oxidation state. FT-EXAFS spectra (Supplementary Fig. 39) reveal a concurrent increase in Co–O coordination, indicating a transformation of the Co environment from tetrahedral to octahedral as lithium is extracted. Wavelet transform analysis (Supplementary Fig. 38) confirms that Co is only coordinated to O or Co neighbors, with no Co–S or Co–S–O coordination detected. ICP–OES analysis of the electrolyte and XPS characterization of the graphite electrode after oxidation confirm the absence of cobalt dissolution or deposition in the electrolyte or on the negative electrode, ruling out the formation of CoSO4, which would easily dissolve in the electrolyte (Supplementary Figs. 40 and 41).

Based on the experimental findings, we propose the following sulfur-mediated oxygen fixation mechanism:

Li6CoO4 + 0.5S → 4Li+ + 4e− + LiCoO2 + 0.5Li2SO4 (2)

Gibbs free energy calculations further support this mechanism (Supplementary Fig. 42). The oxidation pathway of L6CO@S leading to Li2SO4 exhibits a lower Gibbs free energy (ΔG = 593.71 kJ mol−1) compared to direct oxygen evolution from L6CO (ΔG = 1374.52 kJ mol−1). Additionally, the formation of CoSO4 has a higher Gibbs free energy (ΔG = 1052.0 kJ mol−1) than that of Li2SO4, indicating that Li2SO4 formation is thermodynamically favored. This conclusion aligns with our experimental observations, confirming that sulfur preferentially forms Li2SO4 rather than CoSO4.

The proposed reaction mechanism is illustrated schematically in Fig. 3e. Upon charging, S8 molecules may be attacked by reactive oxygen species, leading to S–S bond cleavage and the formation of generating reactive chain-like sulfur intermediates (e.g., [S8]•+, [S]•+, etc.)33. These intermediates gradually react with oxygen species to form SO4 2−, leading to the irreversible formation of the thermodynamically stable Li2SO4. While this mechanism entails a slight compromise in capacity, the oxidation of sulfur to sulfate, rather than the release of gaseous byproducts, markedly suppresses gas-induced side reactions and associate safety concerns. This strategy offers a viable approach to oxygen regulation in high-capacity battery systems through surface oxidation mediators. A comparative analysis between L6CO@S and a range of representative lithium replenishment agents is shown in Fig. 3f and summarized in Supplementary Table 4. Key characteristics, including capacity, decomposition voltage, air stability, and gas evolution, were systematically evaluated. Our proposed L6CO@S demonstrates a well-balanced performance across all four dimensions, highlighting its strong potential for practical application.

Non-destructive Li supply for extended cycle life
The lithium compensation performance of both pristine and sulfur-coated L6CO was quantitatively assessed in Gr||LFP full cells containing 4 wt.% additive. As shown in Fig. 4a, all cells were charged at a constant current of 0.1 C (1 C = 160 mA g−1) in the initial cycle, with a cutoff voltage of 4.3 V to ensure complete decomposition of the lithium supply additive. To determine the real decomposition ratio of L6CO@S when coupled with LiFePO4, additional tests were conducted at 0.1–1 C in Li||LFP+L6CO@S half-cells (Supplementary Fig. 43), showing that the additive is nearly fully oxidized even at 1 C, confirming its compatibility with high-rate formation step. The incorporation of either L6CO or L6CO@S resulted in a pronounced extension of the charge plateau, reflecting the additional lithium input from the additive, and led to a marked increase in cell capacity. The Gr||LFP+4% L6CO@S cell delivered high initial capacities of 205.2 mAh g−1 (charge) and 159.3 mAh g−1 (discharge), approaching the theoretical limit defined by Li||LFP half-cell measurements (Supplementary Fig. 44). This result indicates near-complete restoration of lithium inventory in the full-cell configuration. The rate performance of the LFP full cells also showed improvement with the addition of L6CO and L6CO@S (Supplementary Fig. 45). The long-term cycle stability was further evaluated. After 800 cycles at 0.5 C (1 C = 160 mA g−1), the L6CO@S-based cell retained 92.3% of its initial capacity, markedly outperforming the L6CO counterpart, which exhibited accelerated capacity fade after 500 cycles and retained only 86.8% after 800 cycles (Fig. 4b). Throughout the extended cycling, cells with L6CO@S achieved a high average Coulombic efficiency of 99.993% (Fig. 4c), underscoring the coating’s role in suppressing parasitic side reactions and ensuring stable long-term performance.

In situ electrochemical impedance spectroscopy (EIS) was used to monitor impedance evolution during the initial charging process (Fig. 4d). The Gr||LFP+4% L6CO cell exhibited a progressive increase in impedance as the state of charge (SOC) rose, while the Gr||LFP+4% L6CO@S cell maintained more stable impedance characteristics. Complementary EIS measurements after the 1st, 10th, and 30th cycles revealed a notable increase in resistance for the Gr||LFP+4% L6CO cell (Supplementary Fig. 46), while the Gr||LFP+4% L6CO@S cell showed minimal impedance growth over time. Post-cycling SEM analysis revealed the structural origins of this disparity. The uncoated L6CO electrodes exhibited numerous gas-induced voids and pronounced delamination between the active material layer and the current collector (Fig. 4e and Supplementary Fig. 47), while the L6CO@S electrodes maintained great structural integrity and interfacial adhesion. This morphological contrast can be attributed to gas evolution, as reported in our previous study, where gas generation was found to degrade the electrode structure and lead to the deterioration of electrochemical performance. Moreover, reactive oxygen intermediates chemically attack both electrolyte components and the SEI layer, accelerating impedance growth.

In situ ultrasonic transmission imaging was used to monitor gas evolution during the initial charge process in pouch cells. Leveraging the high acoustic impedance at the gas–liquid interfaces, ultrasound is highly sensitive to the emergence of gas species. As shown in Fig. 4f, the Gr||LFP+4% L6CO cell exhibited gas generation starting at ~60% SOC, with progressive accumulation toward full charge (100% SOC), eventually forming large gas pockets. It is worth noting that, due to the contribution of LFP to overall capacity in the full cell, the onset of gas evolution appears at a higher SOC compared to the ~30% SOC observed in the pure L6CO OEMS test. In contrast, L6CO@S-modified cells maintained consistently high ultrasound transmittance throughout the charging process, showing negligible gas production. These results provide direct visual evidence of the sulfur coating’s effectiveness in suppressing oxygen evolution from L6CO decomposition and subsequent electrolyte oxidation.

To further validate the practical applicability of L6CO@S, its long-term performance was evaluated in Gr||LFP pouch full cells. As shown in Fig. 4g, Gr||LFP pouch cells incorporating L6CO@S delivers an initial discharge capacity that is 12.6% higher than that of control cell, and the improvement becomes more pronounced with extended cycling. After 1400 cycles at 0.5 C (1 C = 160 mA g−1), the capacity enhancement reaches 38.7%, with an overall capacity retention of 91.4%. The corresponding charge/discharge profiles (Supplementary Fig. 48) show minimal polarization growth throughout cycling, indicating well-preserved electrode structure and reaction kinetics. Ultrasonic imaging after 1400 cycles reveals no detectable gas accumulation (Supplementary Fig. 49), and visual inspection shows no swelling of the pouch cell (Supplementary Fig. 50). These collective findings further highlight the critical role of sulfur coating in mitigating gas evolution and enhancing the long-term electrochemical stability. To demonstrate industrial scalability, we further fabricated ampere-hour-scale Gr||LFP pouch cells incorporating L6CO@S additives. As shown in Supplementary Fig. 51, the full-scale cell exhibits great lithium compensation efficacy while maintaining stable cycling performance, confirming seamless integration with commercial manufacturing processes and cell designs. The 1 Ah pouch cell also showed no visible swelling even after 450 cycles (Supplementary Fig. 52). This successful scale-up represents a critical step toward industrial implementation of our upcycling strategy.

The advantages of L6CO@S extend beyond LFP systems to LCO-based full cells. In the Gr||LCO configuration, both initial capacity and cycle life were substantially improved (Supplementary Fig. 53), with capacity retention rising from 83.2% to 93.8% after 300 cycles (Fig. 4h). These results underscore the broad applicability of L6CO@S across different battery systems and highlight its strong potential as a universal lithium-replenishing additive for advanced lithium-ion batteries.

Economic and environmental analysis
A closed-loop lifecycle analysis (LCA) was conducted using the EverBatt 2023 model developed by Argonne National Laboratory to systematically compare the economic and environmental impacts of various recycling approaches. Detailed calculation information is available in the Supplementary Data. Figure 5a–d provides a comprehensive comparison of mainstream battery recycling technologies, along with detailed material flow analyses. Using one ton of spent LCO batteries as an example, approximately 523 kg of SLCO material and 336 kg of negative electrode material can be recovered. Pyrometallurgical recycling involves high-temperature co-smelting of both electrodes, producing metallic alloys such as Co/Cu, while lithium is mostly lost to the slag, resulting in lower overall recovery efficiency and limited product value. In contrast, hydrometallurgical processes use acid-base leaching followed by precipitation to extract lithium and cobalt, ultimately producing about 162 kg of Li2CO3 and 276 kg of CoCO3. Direct regeneration allows the recovery of high-value LCO material. After accounting for process losses and non-repairable LCO, the average yield reaches approximately 421 kg, which is equivalent to 90% of the input SLCO mass. In comparison, our upcycling strategy enables the production of L6CO with higher added value. This process features a streamlined workflow with minimal material loss, avoids extensive acid and base usage, and drastically reduces energy consumption. As a result, approximately 787 kg of L6CO can be obtained.

The energy consumption, greenhouse gas (GHG) emissions, cost-revenue balance, and profitability of various recycling processes were quantitatively assessed. As shown in Fig. 5e, the total recycling costs per kilogram of feedstock are $5.03 (Pyro), $4.79 (Hydro), $4.85 (Direct), and $12.54 (Upcycle). Despite the higher cost of the Upcycle method, 89.8% of this cost is attributed to raw materials (Fig. 5f), rather than energy or infrastructure, indicating high value addition rather than process inefficiency. Revenue analysis (Fig. 5g) shows that the Pyro and Hydro routes yield low-value alloys or salts, while Direct recycling produces regenerated LCO with better marketability. The Upcycle process outperforms all other routes by producing L6CO, a high-performance lithium-replenishment material with great economic potential. Notably, L6CO can be synthesized at an estimated cost of ~$16,000 per ton. Since L6CO is not yet commercially available, we benchmarked its price against LFO (~$69,500 per ton), one of the most representative lithium-rich sacrificial additives already produced at scale. Both L6CO and LFO share an anti-fluorite structure and exhibit suitable decomposition voltages with high irreversible capacities, with L6CO delivering slightly improved electrochemical performance. This pronounced cost–performance advantage highlights the commercial viability and differentiated value of L6CO within next-generation battery manufacturing.

Figure 5h illustrates that the upcycle route delivers the highest profit margin of $42.79 per kilogram of raw material, demonstrating strong economic feasibility. The GHG emissions, shown in Fig. 5i, reveal that the upcycling strategy has the lowest emissions among all methods, at 0.23 kg per kg of feedstock. The rapid Joule heating process also ensures high energy efficiency, consuming just 1.461 MJ per kg of feedstock, considerably less than conventional thermal or wet chemical approaches.

A comprehensive comparison of recycling strategies (Fig. 5j) underscores the superiority of our upcycling approach across multiple dimensions, including profitability, simplicity, recyclability, energy consumption, and GHG emissions. Based on the estimated disposal volume of SLCO batteries in 2024, converting all this waste could yield ~78,700 metric tons of L6CO. This would be sufficient to meet the Li supply demand of ~2.6 million metric tons of LFP materials (assuming a 3 wt.% L6CO additive), exceeding the projected global LFP output of ~2.4 million metric tons in 2024. These figures underscore the potential of this upcycling strategy to alleviate lithium shortages amid the rapid expansion of energy storage and electric vehicle markets. The broader implications of L6CO upcycling method provides a scalable solution that directly transforms waste into performance-enhancing additives, simultaneously improving both sustainability and electrochemical performance throughout the lithium-ion battery value chain.

Methods
Materials and upcycling process
Spent LCO pouch cells were collected from decommissioned mobile phones and laptops, sourced from Apple, Lenovo, ATL, and other manufacturers, with varying degrees of degradation. Spent LiCoO2 (SLCO) batteries were first discharged to 1.0 V (vs. Li+/Li) and then soaked in 1.0 M sodium chloride (NaCl, 99%, Sigma–Aldrich) aqueous solution for 24 h to ensure complete discharge. Disassembly was performed inside an argon-filled glovebox. Recovered electrode foils were rinsed with dimethyl carbonate (DMC, 99%, Sigma–Aldrich) to remove residual electrolyte and lithium salts, then dried under inert atmosphere prior to further processing. The recovered SLCO electrodes were directly employed as working electrodes in a two-electrode electrolysis cell, with a mixed-metal oxide (MMO) mesh as the counter electrode owing to its high oxidation resistance. An aqueous solution composed of 0.5 M sodium sulfate (Na2SO4, 99%, Sigma–Aldrich) with 0.01 M sodium hydroxide (NaOH, 99%, Sigma–Aldrich) was employed as electrolyte. The electrolysis was conducted with a controlled current density of 50 mA cm−2, and the process was terminated once the active material was completely detached. The separated materials were subsequently washed with deionized water, dried at 80 °C for 8 h, and gently ground. The obtained black mass consists of a mixture of SLCO, conductive carbon, and polyvinylidene fluoride (PVDF). While conventional recycling methods necessitate additional steps to eliminate these impurities, our upcycling approach can utilize these impurities as reducing agents during sintering. Hence, they do not require pre-removal. All chemical reagents were used as received.

To convert SLCO to Li6CoO4, 0.291 g of SLCO powder was mixed with 0.232 g of lithium oxide (Li2O, 99.9%, Aladdin) and ground in an agate mortar for 15 min. The powder mixture was then pressed into 14 mm diameter pellets at 6 MPa and held for 5 min, sandwiched between carbon paper, and processed in a Joule heating device (Ultrafast High-Temperature Furnace 2022A, Particle Precision instruments Co. Ltd, jilin Province, China). The temperature was raised to 1200 °C within milliseconds and held for 10 s to yield Li6CoO4. For sulfur coating, the as-prepared Li6CoO4 powder was mixed with elemental sulfur (S, 99.95%, Aladdin) at a mass ratio of Li6CoO4:S = 10:1, hand milled for 15 min, sealed in a hydrothermal reactor under Ar, and heated in a vacuum oven at 160 °C for 6 h to produce L6CO@S.

Battery assembly and electrochemical tests
The LFP and LCO positive electrode slurry was prepared by mixing 80 wt.% active material (LFP from Hubei Wanrun and LCO from Shanshan Corporation), 10 wt.% conductive carbon (Super P, 99.9%, TIMCAL), and 10 wt.% PVDF (99.9%, Aladdin) in N-methyl-2-pyrrolidone (NMP, 99.9%, Aladdin). The slurry was stirred for 20 min in a homogenizer and then coated onto aluminum foil current collectors (15 μm thick, 99.9% purity). For batteries with additional lithium replenishment, L6CO or L6CO@S is directly added to the positive electrode slurry with an additional 4% mass fraction. After drying for 10 h under vacuum at 60 °C, the electrode was cut into 12 mm diameter discs for coin cell assembly. The graphite electrode was prepared similarly, with graphite (Dodochem), Super P, and PVDF mixed at a mass ratio of 9:0.5:0.5, and the slurry was coated onto copper foil (12 μm thick, 99.9% purity). The active materials mass loading of LFP electrodes was ≈7 mg cm−2, and the active materials mass loading of Gr electrodes was 3.65–4 mg cm−2. The N/P ratio was set to 1.1 for the blank batteries and 1.2 for the batteries with lithium supply, respectively. The N/P ratio was calculated on the basis of the reversible capacities of the electrodes.

The CR2025 coin cells were assembled inside an Ar-filled glovebox (25 °C, O2 and H2O levels < 1 ppm). Each cell was filled with 60 μL of electrolyte (1 M LiPF6 dissolved in an ethylene carbonate (EC) and dimethyl carbonate (DMC) mixture (3:7 by volume) with 2 wt.% vinylene carbonate (VC)). Celgard A273 polypropylene membranes (16 μm thickness, 16.5 mm diameter, ~40% porosity, average pore size ~0.039 μm, from Celgard, LLC, USA) were used as separators and were dried at 60 °C for 12 h prior to cell assembly. The positive and negative cases, spring, and spacer, were purchased from Canrd Technology Co. Ltd. Assembly parameters for the 1 Ah pouch cells are summarized in Supplementary Table 5. All coin cells and pouch cells were rested for 8 h to ensure sufficient electrolyte permeation before electrochemical tests. The electrochemical characterizations were carried out at 25 °C using a NEWARE Battery Testing System (MIHW–200–160, Shenzhen, China). Electrochemical impedance spectroscopy (EIS) and cyclic voltammetry (CV) experiments were performed using a GAMRY INTERFACE 1010E electrochemical workstation. EIS measurements were conducted under potentiostatic conditions at open-circuit voltage after the resting period, using an AC amplitude of 10 mV over a 100 kHz–0.1 Hz frequency window with 10 points per decade. The CV scan rate was 0.05 mV s−1. The GITT measurements were conducted with the titration step at 0.1 C for 20 min and a relaxation step of 1 h.

Materials characterization
The composition of SLCO and upcycled materials was determined using inductively coupled plasma optical emission spectroscopy (ICP–OES) (Agilent 5800), while phase compositions were analyzed via X-ray diffraction (XRD) (Bruker D8 Advance) using Cu Kα radiation in the 2θ range of 10–80°. The Fourier transform infrared spectroscopy (FTIR) patterns were tested using the Bruker INVENIO FT-IR analysis platform. Scanning electron microscopy (SEM) (FEI, QUANTA 250 FEG), transmission electron microscopy (TEM) (Tecnai F30), and cryo–TEM (Thermo Fisher Scientific Titan Krios, operated at 300 kV and equipped with a K3 Summit direct electron detection camera, Gatan) were employed to observe the morphology and microstructure of the samples. Energy-dispersive spectroscopy (EDS) mapping was used to analyze the distribution of elements in the particles. The valence state change of sulfur during charging was monitored using X-ray photoelectron spectroscopy (XPS) (Thermo Fisher Scientific Escalab 250Xi) and X-ray absorption spectroscopy (XAS) techniques. XAS data were collected at Shanghai Synchrotron Radiation Facility (SSRF, China). The cells for XAS measurements were disassembled in an argon-filled glovebox, and the electrodes were sealed with Kapton tape and Mylar film to prevent exposure to air. Online differential electrochemical mass spectrometry (OEMS) (QAS 100, Shanghai Linglu Instruments Co., Ltd.) was employed to observe gas evolution during the charging process. Capillary electrophoresis (CE) (P/ACE MDQ capillary electrophoresis system equipped with a UV detector) analysis was conducted to quantitatively determine SO4 2−. The charge electrodes were soaked in deionized water, and the resulting suspension was filtered to obtain the supernatant for analysis. An ultrasonic battery scanner (Wuxi Topsound Technology Co., Ltd.) was used to detect gas generation and distribution within pouch cells.

Theoretical calculations
First-principle calculations were conducted within a plane-wave density-functional theory (DFT) framework. The exchange-correlation potential was treated using the Perdew, Burke, and Ernzerhof (PBE) method of the generalized gradient approximation (GGA). A cutoff energy of 450 eV was adopted for the plane-wave, and the electronic and ionic relaxations were converged to 10−5 eV in solving the Kohn-Sham equations. For all models, a vacuum space of 15 Å was applied along the z-direction to avoid interactions between neighboring images. Sampling of the Brillouin zone was performed using a 2×2×1 k-point grid. All atomic structures were relaxed until the maximum residual force on each atom was below 0.02 eV/Å. Dispersion-corrected DFT calculations were also employed to account for van der Waals interactions. The adsorption strength between the surface and adsorbates was quantified based on the adsorption energy, defined as:

ΔEads = ΔE(adsorb/surf) − ΔE(surf) − ΔE(adsorb)

Where ΔE(adsorb/surf), ΔE(surf) and ΔE(adsorb) correspond to the calculated total energies of the substrate with adsorbate(s), the clean substrate, and the isolated adsorbate, respectively.

All optimized atomic configurations used in the above calculations are provided in Supplementary Data 1.

Economic analysis of different battery recycling technologies
EverBatt developed by Argonne National Laboratory was used to evaluate the cost, profit, energy consumption, and GHG emissions of pyro-, hydro-, direct recycling, and our upcycling method. A total of 10,000 tonnes of end-of-life LCO batteries were input for preprocessing, with the separated black mass processed via various recycling methods. A detailed assessment of raw material costs and potential product revenues is provided in Supplementary Tables 6–11. Due to the high lithium content in L6CO, the raw material costs, particularly for lithium salts, are higher than in other recycling methods. The market price for L6CO was estimated based on the price of LFO. Additionally, the GHG emissions and total energy consumption of the different processes were evaluated, as detailed in Supplementary Table 12.

Data availability
Source data are provided with this paper.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This is an interesting study that advances both spent battery recycling and practical prelithiation strategy. While waste\nLiCoO2 batteries represent a valuable resource due to their high cobalt and lithium content, directly regenerating degraded\ncathodes remains a complex process with multiple technical and chemical challenges. The strategy proposed in this work\nfor converting waste LiCoO2 to Li6CoO4 is simple, rapid, and highly tolerant to feedstock variability. Additionally, the simple\nsulfur coating provides an effective solution to address challenges related to stability and oxygen release. The manuscript is\nwell-prepared and logically structured. The authors present convincing electrochemical demonstrations and thorough\ntechno-economic assessments, highlighting the industrial relevance of their approach. However, several points in the\nmanuscript require further clarification and a more detailed discussion. I recommend acceptance after minor revisions, as\noutlined below.\n1. The manuscript claims that converting SLCO into Li6CoO4 provides a more “universal” recycling route compared with\ndirectly regenerating it into LCO. Could the authors clarify why the Li6CoO4 pathway is more universal than the conventional\nLCO regeneration approach?\n2. Different sulfur contents are expected to significantly influence performance. How was the sulfur coating ratio determined?\nIs 10 wt.% an optimized value, or could other loadings be more effective?\n3. Considering that sulfur has poor electronic conductivity and may affect electron transport, does the sulfur coating influence\nthe reaction kinetics of Li6CoO4?\n4. During long-term air storage, does the sulfur layer undergo spontaneous oxidation? If degradation occurs during storage\nor transportation, it may affect industrial applicability.\n5. Some experimental conditions are not fully described, which may affect reproducibility. Please provide the detailed\nexperimental steps for the electrode separation process.\n6. Since SLCO was used without additional pretreatment, how do impurities such as residual PVDF or aluminum influence\nboth the conversion process and the subsequent cycling performance of the batteries?\n7. Considering that Li2O is relatively expensive and that in many cathode studies, lithium precursors are commonly chosen\nfrom Li2CO3 or LiOH, why was Li2O chosen as the lithium source, instead of more common and lower-cost options such as\nLiOH or Li2CO3?\n8. In the economic analysis, the authors benchmark the price of Li6CoO4 against Li5FeO4, which is relatively expensive.\nThis may inflate the calculated economic value. Could the authors provide a justification for this benchmark?", "reviewer_2": "The authors developed a FJH method to upcycle spent LCO and demonstrated the L6CO as a high-capacity sacrificial\nadditive. Furthermore, the sulfur coating was used to improve the stability of L6CO, which seems promising for different full\ncells. The characterizations, including spectroscopy, microscopy and electrochemistry, are systematic, and the results are\npromising. Overall, this work demonstrates high innovation, but several issues in the manuscript require further\nimprovement.\n1. The author claimed the upcycling conversion can be achieved within 10 seconds, but the sulfur coating post-treatment\nstep is rather time-consuming (6 hours), making the whole process less promising. I wonder if the post-treatment can be\nachieved using the same FJH method, maybe a stepwise FJH approach to highlight the advantages of the proposed\nmethod.\n2. For spent cathode material, it includes electrolyte, CEI, as well as binders, which has a high content of fluorine. I wonder\nits reaction pathway during flash upcycling method, its potential environmental effects and how to solve them.\n3. Another problem related to spent cathode is the impurities, such as Al and Cu, which can negatively affect the\nelectrochemical performance as demonstrated in the previous FJH work. I wonder the contents of these impurities and if\nthese impurities can affect the performance of L6CO.\n4. “Based on a comprehensive assessment of reaction kinetics and energy efficiency, 1200°C for 10 seconds was identified\nas the optimal synthesis condition (Supplementary Fig. 7).” Can the authors explain more details about reaction kinetics and\nenergy efficiency, I cannot get the rationale.\n5. “The LFP+4% L6CO@S||Gr cell delivered outstanding initial capacities of 205.2 mAh g−1 (charge) and 159.3 mAh g−1\n(discharge)”. I wonder if there are >10% extra Li stored at the anode side, will the first charging cycle cause lithium metal\ndeposition at the anode side, since the N/P ratio is 1.1?\n6. S is an insulator and substantial amount of conductive carbon is required in Li-S batteries. Here, S coating is used, and I\nwonder why it does not harm the specific capacity and rate performance.\n7. I wonder by adding extra L6CO@S, what the increasing ratios of specific capacity in full pouch cells are?\n8. More details about the flash upcycling method should be provided, including reactant mass, reactor dimensions and so\non. Since the heat and mass transfer are directly related to reaction condition, the upscaling ability should be demonstrated.\n9. The LCA only considers the formation of L6CO, but the experiment parts focus more on L6CO part, I wonder if the whole\nprocedure can be considered for the upcycling method. And how is the price of L6CO estimated, since it should be highly\nperformance and quality-determined?\n10. Recent spent batteries have diverse chemistry, like NMC, NCA or their mixture within spent cathodes, how can the flash\nupcycling method be used in these spent cathodes?", "reviewer_3": "In this study, the authors propose a novel upcycling method for converting spent LiCoO₂ into Li₆CoO₄ via flash Joule heating,\nachieving complete lithium and cobalt recovery within 10 seconds. The synthesized Li₆CoO₄ is further sulfur-coated to\nmitigate air sensitivity and oxygen release during delithiation, ultimately serving as a high-capacity sacrificial additive to\noffset active lithium loss. The work demonstrates promising electrochemical performance when integrated into\nLiFePO₄||Graphite pouch cells, achieving 91.4% capacity retention over 1400 cycles. The manuscript is commendable for its\nwell-integrated narrative that connects material design, mechanistic understanding, and electrochemical validation.\nHowever, certain aspects require further clarification, particularly regarding material synthesis, the redox mechanism of\noxygen anions and sulfur, and battery cycling performance. Overall, this is a meaningful work and minor revisions are\nrecommended:\n1. The waste LiCoO₂ particles exhibit Co₃O₄ on the surface, with the bulk material remaining LiCoO₂. How is full conversion\nof LiCoO₂ to Li₆CoO₄ achieved via flash Joule heating, and how is stoichiometry of LixCoyO₂ accurately determined? Why\ncan’t CoO be fully converted to Li₆CoO₄ by Joule heating?\n2. While the manuscript states that the modified Li₆CoO₄@S exhibits excellent rate performance, has this performance been\nevaluated when coupled with LiFePO₄? Specifically, how does the rate performance vary at different initial cycle currents?\nWhat is the observed decomposition rate under these conditions?\n3. How is the sulfur coating uniformly applied to Li₆CoO₄? Could the authors provide further details on the coating process?\nGiven sulfur's poor conductivity, how does the sulfur coating impact the impedance of the material and its decomposition rate\nduring cycling? Furthermore, how is the sulfur loading at 160°C controlled to ensure an amorphous state?\n4. At 40% and 60% state of charge (SOC), the S K-edge XANES spectra show peaks for sulfur and Li₂SO₄, along with an\nadditional smaller peak. What compound is represented by this smaller peak? Are there any intermediate phases that could\ninfluence the battery system? Moreover, aside from soft X-ray analysis, are there other methods used to confirm the\nformation of Li₂SO₄ and provide direct evidence of the reaction between sulfur and reactive oxygen to generate [S₈]⁺ and\n[S₅]⁺?\n5. The manuscript suggests that the introduction of the prelithiation agent enhances both the discharge capacity and cycling\nstability of the LiFePO₄||Graphite full cell. However, the Li₆CoO₄ prelithiation agent primarily affects the first cycle and is\nunlikely to have a significant impact on subsequent cycles. Please elaborate on the mechanism by which the prelithiation\nagent enhances cycling performance over subsequent cycles. Additionally, since LiFePO₄||Graphite cells are mature\ncommercial systems with good cycling stability even without prelithiation, the significant capacity degradation observed in\nFigures 4b, 4g, and 4h requires further explanation.\n6. Figure 3i is cited as comparing various prelithiation agents in terms of air stability and gas evolution, but this figure does\nnot appear in the manuscript. Does the reference actually pertain to Figure 3f? Please clarify."}, "author_response": "Responses to Reviewer Comments\nTable of Contents\n1. Response to Reviewer #1 .................................................................................................................... 2\n2. Response to Reviewer #2 .................................................................................................................. 15\n3. Response to Reviewer #3 .................................................................................................................. 33\n1 / 45\nResponse to Reviewer #1\nComment:\nThis is an interesting study that advances both spent battery recycling and practical prelithiation strategy.\nWhile waste LiCoO batteries represent a valuable resource due to their high cobalt and lithium content,\n2\ndirectly regenerating degraded cathodes remains a complex process with multiple technical and\nchemical challenges. The strategy proposed in this work for converting waste LiCoO to Li CoO is\n2 6 4\nsimple, rapid, and highly tolerant to feedstock variability. Additionally, the simple sulfur coating\nprovides an effective solution to address challenges related to stability and oxygen release. The\nmanuscript is well-prepared and logically structured. The authors present convincing electrochemical\ndemonstrations and thorough techno-economic assessments, highlighting the industrial relevance of\ntheir approach. However, several points in the manuscript require further clarification and a more\ndetailed discussion. I recommend acceptance after minor revisions, as outlined below:\nResponse:\nThank you for your thoughtful and constructive evaluation of our manuscript. We are very pleased\nthat you recognize the novelty and industrial relevance of our upcycling strategy. In response to your\nfeedback, we have made several targeted revisions to strengthen the manuscript:\na) We expanded the discussion to clarify why the Li CoO pathway is considered more universal\n6 4\nthan direct LCO regeneration, emphasizing its tolerance to diverse degradation states and its functional\nversatility as a sacrificial lithium donor.\nb) The sulfur coating section has been revised with additional discussion on sulfur loading, its\noptimization, and its influence on reaction kinetics.\nc) The Methods and Supplementary Information now provide a detailed step-by-step description\nof the electrode separation procedure to ensure reproducibility.\nd) Additional explanation has been added on the influence of impurities such as residual PVDF\nand Al. Supporting FTIR and ICP data have been included.\ne) The rationale for selecting Li O as the lithium source has been expanded, and the economic\n2\nanalysis has been revised to justify the use of Li FeO as a rational benchmark\n5 4\nIn the following, we provide point-by-point responses to your comments. The content cited from\nthe manuscript and Supplementary files is marked in blue in this document, while the corresponding\nrevisions made in both the manuscript and Supplementary files are highlighted in yellow. We sincerely\nhope that the revised version addresses all your justified concerns and meets the expectations for a\ndecent Nature Communications publication.\n2 / 45\nComment:\n1. The manuscript claims that converting SLCO into Li CoO provides a more “universal” recycling\n6 4\nroute compared with directly regenerating it into LCO. Could the authors clarify why the Li CoO\n6 4\npathway is more universal than the conventional LCO regeneration approach?\nResponse:\nWe appreciate your insightful question regarding the universality of the Li CoO pathway. The\n6 4\ndistinction lies in the inherent limitations of conventional LCO regeneration compared with the broader\napplicability of the Li CoO strategy.\n6 4\nDirect regeneration of degraded LCO requires strict control over the precursor’s chemistry,\nstructure, and degradation state. In practice, cathode waste from different sources exhibits highly\nvariable failure modes, including lithium depletion, lattice collapse, cobalt dissolution, cation mixing,\nand surface contamination by Li CO , LiOH, residual carbon, or PVDF.1-3 These defects make it\n2 3\ndifficult to fully restore the layered structure. Even if XRD patterns appear to show partial restoration\nof the layered phase, the long-term electrochemical performance often remains unsatisfactory.\nMoreover, direct LCO regeneration often requires multi-step treatments such as high-temperature solid-\nstate synthesis, molten-salt recrystallization, or hydrothermal processes, and the exact route must be\ntailored to the degradation level of the precursor.4,5 This makes direct LCO regeneration difficult to\nimplement as a general recycling route for heterogeneous industrial waste streams.\nIn contrast, the Li CoO pathway relies on reconstructing a new antifluorite structure rather than\n6 4\npreserving the parent layered framework. The conversion is governed primarily by the lithium-to-cobalt\nstoichiometry and is largely independent of the precursor’s degradation state. Whether the feedstock is\nmildly aged or severely failed, Li CoO can be rapidly obtained through flash Joule heating (FJH) at\n6 4\n~1200°C for 10 s, directly utilizing residual carbon and PVDF binders as reductants. This simple, one-\nstep process eliminates the need for elaborate pretreatments or sorting, thereby improving scalability\nand tolerance to feedstock variability. Furthermore, unlike regenerated LCO, Li CoO is not intended\n6 4\nto serve as a cycling cathode itself. Instead, it functions as a high-capacity, sacrificial lithium donor that\ncan enhance the performance of diverse commercial cathodes (e.g., LFP, LCO, and NCM). This\nfunctional versatility effectively redirects the value of spent LCO into a broader range of practical\napplications.\nFor these reasons, the Li CoO route can be considered a more universal and broadly applicable\n6 4\nrecycling strategy compared with conventional LCO regeneration, which is typically constrained to\nnarrow feedstock conditions and specific end uses.\nThis revised section has been added to the revised manuscript (Page 3, lines 10–22, highlighted):\n“This approach circumvents the need for complex repair of the layered framework required for\n3 / 45\nstable cycling, but instead reconstructs an antifluorite structure that irreversibly releases active lithium\nto enhance energy density and extend battery lifespan. L6CO does not function as a cycling cathode but\nbecomes electrochemically inactive after initial decomposition19, thus avoiding concerns over long-\nterm stability.”\n“This method also demonstrates high universality, enabling efficient conversion of SLCO across\nvarious degradation states while achieving near-complete recovery of both lithium and cobalt.”\nComment:\n2. Different sulfur contents are expected to significantly influence performance. How was the sulfur\ncoating ratio determined? Is 10 wt.% an optimized value, or could other loadings be more effective?\nResponse:\nWe thank you for raising this important question. The sulfur loading in our work was determined\nby combining stoichiometric calculation with experimental verification. The design of the sulfur\ncoating was initially guided by the following reaction:\nLi CoO + 0.5 S → 4 Li+ + 4 e−+ LiCoO + 0.5 Li SO\n6 4 2 2 4\nIn this pathway, sulfur acts as an oxygen scavenger that converts lattice oxygen intermediates into\nstable sulfate. According to this equation, the required sulfur amount corresponds to ~10 wt.% relative\nto Li CoO . We also considered another possible pathway:\n6 4\nLi CoO + S → 6 Li+ + 6 e−+ CoSO\n6 4 4\nThis would lead to the formation of cobalt sulfate, with a theoretical capacity of ~800 mAh g−1.\nHowever, Gibbs free energy calculations showed that this reaction is thermodynamically unfavorable\ncompared with the formation of Li SO . Furthermore, in our experiments no signals of CoSO were\n2 4 4\ndetected, which strongly supports the first reaction pathway. Importantly, the measured electrochemical\ncapacity is consistent with the stoichiometry of the first reaction rather than the second.\nTo further evaluate the optimal sulfur ratio, we prepared L6CO@S samples with different sulfur\nloadings (5, 10, and 20 wt.%). The results clearly confirmed the trade-off between oxygen suppression\nand charge capacity (Figures. R1 and R2). With 5 wt.% sulfur loading, the coating was insufficient to\nfully stabilize oxygen intermediates. OEMS measurements revealed detectable oxygen release, and the\ninterfacial stability remained poor. At 20 wt.%, a thicker sulfur layer achieved complete oxygen\nsuppression，but the excess sulfur remained electrochemically inactive, lowering the overall specific\ncapacity and slightly increasing polarization. At 10 wt.%, oxygen evolution was already effectively\nsuppressed, while maintaining balanced kinetics and capacity.\n4 / 45\nFigure R1. Initial charge curves of L6CO@S with different sulfur contents (0–20 wt.%) at 0.05\nC.\nFigure R2. In-situ OEMS spectra of L6CO@S with different sulfur contents (0–20 wt.%).\nTaken together, both the theoretical calculation and the experimental validation converge on 10\nwt.% sulfur as the optimal loading, offering the best compromise among air stability, capacity, and the\neffectiveness of oxygen suppression.\nComment:\n3. Considering that sulfur has poor electronic conductivity and may affect electron transport, does the\n5 / 45\nsulfur coating influence the reaction kinetics of Li CoO ?\n6 4\nResponse:\nThank you for this important question. Under the coating protocol used in this work (10 wt.% S),\nthe deposited sulfur forms a thin, conformal shell that does not sever the electronic percolation network\nbetween L6CO particles and the conductive additive; consequently, the coating does not produce a\nsubstantive kinetic bottleneck under practical operating conditions.\nImpedance measurements (Figure R3) further confirm that the charge-transfer resistance of\nL6CO@S is comparable to that of pristine L6CO, indicating that the sulfur layer does not significantly\nalter the electrode resistance. Moreover, given the small additive loading (4 wt.% of the total cathode,\nwith sulfur accounting for only ~0.4 wt.%), the coating exerts a negligible influence on the overall\nimpedance of the composite electrode.\nAs shown in Figure R1, a modest increase in polarization is observed as sulfur loading rises, but\nthe accompanying capacity loss is not solely a conductivity-driven effect. The dominant cause of the\nreduced charge capacity is a change in the oxygen-scavenging reaction pathways. To evaluate the direct\ninfluence of the coating on transport and kinetics, we performed galvanostatic intermittent titration\ntechnique (GITT) measurements and derived Li⁺ diffusion coefficients (D ) across the state-of-charge\nLi\nrange (Figure R4). The D values of pristine L6CO and L6CO@S (10 wt.% S) remain in the same order\nLi\nof magnitude, indicating that lithium-ion transport is essentially preserved after sulfur coating.\nFigure R3. a, Electrochemical impedance spectroscopy (EIS) spectra of L6CO||Li and L6CO@S||Li\nhalf cells. b, EIS spectra of LFP||Li, LFP+4 % L6CO||Li, and LFP+4 % L6CO@S||Li half cells.\n6 / 45\nFigure R4. a,b, Voltage profiles of GITT measurements for L6CO and L6CO@S. c, Corresponding Li⁺\ndiffusion coefficients derived from the GITT data.\nThe GITT data have been added to the revised manuscript and Supplementary Information (Page\n8, lines 4–7), along with experimental methods (Page 18, lines 17–19, highlighted):\n“Moreover, GITT measurements reveal that the Li⁺ diffusion coefficients (D ) of pristine L6CO\nLi\nand L6CO@S remain in the same order of magnitude before and after coating (Supplementary Fig. 26),\nindicating that lithium-ion transport is not significantly hindered by the sulfur layer.”\n“The GITT measurements were conducted with the titration step at 0.1C for 20 min and a\nrelaxation step of 1 h.”\nComment:\n4. During long-term air storage, does the sulfur layer undergo spontaneous oxidation? If degradation\noccurs during storage or transportation, it may affect industrial applicability.\nResponse:\nWe appreciate your concern regarding the air stability of the sulfur coating during storage and\ntransportation. To clarify, we conducted air-exposure tests on L6CO@S and measured the surface\nchemistry by XPS after 3 and 7 days of storage. The S 2p spectra show that sulfur remains dominated\nby elemental features during storage (Figure R5), no discernible increase in the high-binding-energy\ncomponents characteristic of sulfate formation. These results indicate that the sulfur coating is stable\n7 / 45\nunder normal ambient storage. The oxidation to sulfate in our system requires both electrochemical\ndriving force and highly reactive oxygen species (e.g., radical oxygen generated during charging\nprocess).\nFigure R5. S 2p XPS spectra of L6CO@S after air exposure for different storage durations.\nThe above XPS data have been added to the revised manuscript and Supplementary Information\n(Page 11, lines 5-10):\n“In addition, we evaluated the air stability of the sulfur coating. S 2p XPS spectra of L6CO@S\nstored in air for 3 and 7 days showed no evidence of sulfate formation (Supplementary Figure 36),\nindicating that the sulfur does not undergo spontaneous oxidation in ambient conditions. The oxidation\nto sulfate in our system requires both electrochemical driving force and the presence of highly reactive\noxygen species.”\nComment:\n5. Some experimental conditions are not fully described, which may affect reproducibility. Please\nprovide the detailed experimental steps for the electrode separation process.\nResponse:\nWe appreciate your suggestion on providing more methodological detail to improve reproducibility.\nIn response, we have included the following detailed experimental methods in the revised manuscript (Pages\n17–18, highlighted).\na) Disassembly and pre-treatment of spent pouch cells:\n8 / 45\n“Spent LiCoO (SLCO) batteries were first discharged to 1.0 V and then soaked in 1.0 M NaCl\n2\naqueous solution for 24 h to ensure complete discharge. Disassembly was performed inside an argon-\nfilled glovebox. Recovered electrode foils were rinsed with dimethyl carbonate (DMC) to remove\nresidual electrolyte and lithium salts, then dried under inert atmosphere prior to further processing.”\nb) Water electrolysis-induced separation (WES) process:\n“The recovered cathodes were directly employed as working electrodes in a two-electrode\nelectrolysis cell, with a mixed-metal oxide (MMO) mesh as the counter electrode owing to its high\noxidation resistance. An aqueous solution composed of 0.5 M Na SO with 0.01 M NaOH was\n2 4\nemployed as electrolyte. The electrolysis was conducted with a controlled current density of 50 mA\ncm−2, and the process was terminated once the active material was completely detached. The separated\nmaterials were subsequently washed with deionized water, dried at 80°C for 8 h, and gently ground.\nThe obtained black mass consists of a mixture of SLCO, conductive carbon, and polyvinylidene fluoride\n(PVDF).”\nc) Flash Joule heating (FJH) upcycling and sulfur coating process:\n“To convert SLCO to Li CoO , 0.291 g of spent cathode powder was mixed with 0.235 g of Li O\n6 4 2\nand ground in an agate mortar for 15 min. The powder mixture was then pressed into 14 mm diameter\npellets, sandwiched between carbon paper, and processed in a Joule heating device (Ultrafast High-\nTemperature Furnace 2022 A, Particle Precision instruments Co. Ltd, jilin Province, China). The\ntemperature was raised to 1200 °C within milliseconds and held for 10 s to yield Li CoO . For sulfur\n6 4\ncoating, the as-prepared Li CoO powder was mixed with elemental sulfur at a mass ratio of\n6 4\nLi CoO :S=10:1, hand milled for 15 min, sealed in a hydrothermal reactor under Ar, and heated in a\n6 4\nvacuum oven at 160°C for 6 h to produce L6CO@S.”\nComment:\n6. Since SLCO was used without additional pretreatment, how do impurities such as residual PVDF or\naluminum influence both the conversion process and the subsequent cycling performance of the\nbatteries?\nResponse:\nThank you for this important question. In the flash Joule heating (FJH) process, the polymer binder\nPVDF in the SLCO electrodes serves as an effective reducing agent. To demonstrate this, we performed\na control experiment using commercial LCO as the starting material, mixed with PVDF and Li O.\nThe\ncarbon black was added so the PVDF served as the sole reductant. After subjecting the mixture to\nidentical FJH conditions, XRD results confirmed that the LCO was completely converted into L6CO\n(Figure R6). This clearly indicates that PVDF decomposes under FJH and directly participates in the\n9 / 45\nreduction chemistry. Consistent with this, FTIR spectra of SLCO before FJH exhibit multiple absorption\npeaks in the 900–1400 cm−1 range corresponding to PVDF6,7, whereas these peaks almost completely\ndisappear after FJH (Figure R7), confirming that the binder fully decomposes and does not remain in\nthe final product.\nFigure R6. XRD patterns of products obtained from LiCoO when PVDF was used as the sole reductant\n2\nunder FJH (1200°C, 10 s).\nFigure R7. FTIR spectra of SLCO before and after FJH conversion to L6CO.\nRegarding aluminum, our water electrolysis separation (WES) procedure effectively minimizes Al\ncontamination from current collectors. As shown in Figure R8, the aluminum foils were completely\nseparated from the active material without fracture. ICP analysis further confirmed that the residual Al\ncontent in the recovered SLCO powder was only 1.8 ppm, and no Cu impurity was detected (Figure\nR9). Moreover, L6CO is not used as a cathode material itself but as a lithium supplement, with an\n an\naddition amount of only ~4 wt.% relative to the cathode. Therefore, even if trace Al impurities were\npresent, the resulting Al concentration in the overall electrode would be far below levels that could\n10 / 45\ninfluence battery performance.\nFigure R8. a, Photograph of the WES setup. b, Optical image of the SLCO electrode before WES. c,\nRecovered Al foil and separated SLCO cathode powder after WES.\nFigure R9. ICP results showing the concentrations of Al and Cu impurities in SLCO.\nIn conclusion, PVDF serves as a reductant during the FJH process, facilitating the conversion of\nSLCO into L6CO, while residual Al impurity remains at extremely low levels in SLCO (molar fraction\n< 1%), which is insufficient to influence the electrochemical behavior of L6CO as a lithium supplement.\nThe FTIR and ICP data are included in the revised manuscript to clarify these points. (Page 4, lines 15–\n17 and Page 6, lines 20–24, highlighted):\n“ICP analysis further confirmed that the residual Al content in the SLCO powder was only 1.8 mg\nL⁻¹, and no Cu impurity was detected (Supplementary Fig 2).”\n“FTIR of SLCO displays characteristic absorptions peaks corresponding to PVDF (900–1400\ncm⁻¹), which almost completely disappear after FJH (Supplementary Fig. 13), confirming its full\ndecomposition. Owing to the brief reaction duration, the released fluorinated volatiles are transient and\nconcentrated in the off-gas stream, which makes them amenable to engineering control.”\nComment:\n7. Considering that Li O is relatively expensive and that in many cathode studies, lithium precursors\n2\nare commonly chosen from Li CO or LiOH, why was Li O chosen as the lithium source, instead of\n2 3 2\nmore common and lower-cost options such as LiOH or Li CO ?\n2 3\nResponse:\nThank you for this insightful question. In the context of ultrafast, high-temperature FJH, it is\ncrucial to employ a lithium precursor with sufficiently high chemical reactivity to drive a complete\nsolid-state transformation within seconds. Li O fulfills this requirement most effectively, enabling clean\n2\nand phase-pure formation of L6CO under FJH conditions. Consistent with previous studies, lithium-\nrich compounds such as Li FeO have also been synthesized using Li O as the lithium source8,9. In\n5 4 2\ncontrast, when Li CO or LiOH is used under the same rapid heating, incomplete decomposition often\n2 3\nleaves residual phases or carbonate byproducts, which compromise both the phase purity and the\nelectrochemical stability of the final material.\nFrom an industrial perspective, Li O can be readily produced from inexpensive and abundant\n2 3\nfeasible. At the same time, we acknowledge that other precursors can also be employed if the processing\nconditions are appropriately modified. For example, by replacing Li O with LiOH and applying\n2\nconventional furnace treatment (800°C, 12 h under Ar), we were able to obtain phase-pure L6CO as\nconfirmed by XRD (Figure R10). This approach demonstrates that alternative lithium sources are viable,\nalthough it requires prolonged heating and therefore sacrifices the time and energy advantages of FJH.\n12 / 45\nFigure R10. XRD pattern of L6CO obtained using LiOH as the lithium source via conventional tube-\nfurnace sintering (800°C, 12 h, Ar).\nComment:\n8. In the economic analysis, the authors benchmark the price of Li CoO against Li FeO , which is\n6 4 5 4\nrelatively expensive. This may inflate the calculated economic value. Could the authors provide a\njustification for this benchmark?\nResponse:\nWe agree with the Reviewer that the choice of an appropriate benchmark is critical for a fair\neconomic assessment. Since L6CO is not yet a commercial product, no direct market price is available.\nWe therefore selected Li FeO as the comparator because it is one of the most representative lithium-\n5 4\nrich sacrificial additives reported to date and is already available at scale9.\nAlthough lithium oxalate (Li C O ) has been investigated as a lower-cost lithium source, its high\n2 2 4\ndecomposition potential (~4.5 V) makes it impractical for real applications.\nBy contrast, both L6CO and Li FeO exhibit suitable decomposition voltages and high irreversible\n5 4\ncapacities, and they share a similar anti-fluorite crystal structure. In fact, L6CO delivers slightly\nsuperior electrochemical performance compared to Li FeO , making the latter a reasonable and\n5 4\nconservative reference point for economic analysis. Using Li FeO as a benchmark therefore does not\n5 4\noverstate the economic value of L6CO, but rather provides a fair and functionally relevant basis for\ncomparison. We have clarified this rationale in the revised manuscript (Page 16, lines 9–16, highlighted):\n“Notably, Li CoO can be synthesized at an estimated cost of ~$16,000 per ton. Since Li CoO is\n6 4 6 4\nnot yet commercially available, we benchmarked its price against Li FeO (~$69,500 per ton), one of\n5 4\nthe most representative lithium-rich sacrificial additives already produced at scale. Both Li CoO and\n6 4\nLi FeO share an anti-fluorite structure and exhibit suitable decomposition voltages with high\n5 4\nirreversible capacities, with Li CoO delivering slightly superior electrochemical performance. This\n6 4\n13 / 45\npronounced cost–performance advantage highlights the commercial viability and differentiated value\nof Li CoO within next-generation battery manufacturing.”\n6 4\nReferences:\n1. Lyu, Y. et al. An Overview on the Advances of LiCoO Cathodes for Lithium‐Ion Batteries. Adv. Energy Mater.\n2\n11, 2000982 (2020).\n2. Jia, K. et al. Degradation Mechanisms of Electrodes Promotes Direct Regeneration of Spent Li‐Ion Batteries:\nA Review. Adv. Mater. 36, 2313273 (2024).\n3. Meng, Z., Ma, X., Hou, J., Zheng, Y. & Wang, Y. Impurity Impacts of Recycling NMC Cathodes. Adv. Energy\nMater. (2025).\n4. Wu, X. et al. Toward Circular Energy: Exploring Direct Regeneration for Lithium‐Ion Battery Sustainability.\nAdv. Mater. 36 (2024).\n5. Ma, X. et al. The evolution of lithium-ion battery recycling. Nat. Rev. Clean Technol. 1, 75–94 (2025).\n6. Cai, X., Lei, T., Sun, D. & Lin, L. A critical analysis of the α, β and γ phases in poly(vinylidene fluoride) using\nFTIR. RSC Adv. 7, 15382-15389 (2017).\n7. Meng, N. et al. Ultrahigh β-phase content poly(vinylidene fluoride) with relaxor-like ferroelectricity for high\nenergy density capacitors. Nat. Commun. 10 (2019).\n8. Zhu, B. et al. Understanding the Air‐Exposure Degradation Chemistry of the Sacrificial Cathode Additive\nLi5FeO4 for Li‐Ion Batteries. Adv. Funct. Mater. 34, 2315010 (2024).\n9. Liu, C. et al. Air-stable Li FeO additive enabled by carbon coating for energy-dense lithium-ion batteries. Nat.\n5 4\nCommun. 16 (2025).\n14 / 45\nResponse to Reviewer #2\nComment:\nThe authors developed a FJH method to upcycle spent LCO and demonstrated the L6CO as a high-\ncapacity sacrificial additive. Furthermore, the sulfur coating was used to improve the stability of L6CO,\nwhich seems promising for different full cells. The characterizations, including spectroscopy,\nmicroscopy and electrochemistry, are systematic, and the results are promising. Overall, this work\ndemonstrates high innovation, but several issues in the manuscript require further improvement.\nResponse:\nThank you for the careful and constructive assessment of our work. We greatly appreciate your\nrecognition of the significance of our work and the innovation of our upcycling approach. In response\nto your detailed comments, we have made a series of focused revisions and additions to improve the\nclarity and completeness of the manuscript:\na) We experimentally examined the possibility of stepwise FJH for sulfur coating and explained\nits limitations. Process-intensification alternatives were also discussed.\nb) The fate of fluorine-containing components (PVDF, CEI, residual electrolyte) during FJH was\ninvestigated, supported by control experiments and FTIR analysis. Environmental implications and\nmitigation strategies were discussed. Metallic impurities (Al, Cu) were also quantified by ICP.\nd) The rationale for selecting 1200°C/10 s as the optimal FJH condition was expanded, supported\nby XRD results and kinetic analysis.\ne) N/P ratio design and lithium plating risk were explicitly analyzed.\nf) We presented measurements quantifying the effect of the sulfur coating on the kinetics of L6CO.\ng) Full-pouch-cell comparisons were added, and more detailed experimental procedures were\nprovided. The scalability of the FJH process was clarified.\nh) The scope of LCA and economic assessment was clarified, and the rationale for benchmarking\nagainst Li FeO was provided.\n5 4\ni) The potential extension of FJH to NCM cathodes was explored experimentally, with preliminary\nresults and future directions included.\nFor ease of cross-checking, the content cited from the manuscript or Supplementary Information\nis marked in blue in this response document, while the exact revisions we implemented in the revised\nmanuscript and Supplementary files are highlighted in yellow. We sincerely hope that these revisions\naddress the Reviewer’s concerns and substantially strengthen the manuscript.\n15 / 45\nComment:\n1. The author claimed the upcycling conversion can be achieved within 10 seconds, but the sulfur\ncoating post-treatment step is rather time-consuming (6 hours), making the whole process less\npromising. I wonder if the post-treatment can be achieved using the same FJH method, maybe a\nstepwise FJH approach to highlight the advantages of the proposed method.\nResponse:\nWe appreciate your constructive suggestion to explore a stepwise FJH route for sulfur coating. We\nconsidered this option experimentally, but found that FJH is not well suited to produce a thin, conformal\nsulfur coating for the reasons explained below.\nFlash Joule heating is a current-driven, highly transient heating method that requires the sample to\nbe electrically coupled (typically sandwiched between conductive carbon paper) and subjected to a\ntransient current pulse in our setup. The minimum controllable temperature in our FJH system is 400 ℃,\nwhich is close to the boiling point of elemental sulfur, and the relatively large chamber volume prevents\nlocal sealing. Such elevated, short-duration heating produces rapid sulfur volatilization and prevents\nuniform re-deposition on particle surfaces. To verify the feasibility, we uniformly mixed L6CO and\nelemental sulfur and subjected the mixture to FJH treatment under various conditions (400 °C, 1–10 s).\nThe obtained samples showed no noticeable color change (Figure R1), whereas successful sulfur-coated\nproducts typically exhibit a dark-green appearance. Consistently, XPS analysis of the sample treated at\n400 °C for 10 s (Figure R2) revealed negligible surface sulfur, indicating that sulfur was largely\nvolatilized or agglomerated rather than conformally deposited under these transient conditions.\nFigure R1. Optical photographs of L6CO samples before and after attempted sulfur coating via FJH\nat 400 °C for different heating durations.\n16 / 45\nFigure R2. S 2p XPS spectra of L6CO and sulfur compound treated by FJH at 400°C for 10 s in an\nattempt to form a sulfur coating.\nFor these practical and physicochemical reasons, we developed the low-temperature sealed-vessel\nprotocol reported in the manuscript. Heating L6CO together with elemental sulfur at 160°C in a sealed\nreaction vessel enables molten sulfur to wet particle surfaces and redistribute slowly, producing a\nuniform, conformal sulfur layer without excessive volatilization. To reconcile throughput with coating\nquality, we also optimized the post-treatment to shorten its duration while preserving coating uniformity.\nWe reduced the required processing time from 6 h to as little as 3 h while retaining a thin, homogeneous\ncoating as verified by cryo-TEM (Figure R3).\nFigure R3. Cryo-TEM image of L6CO@S (5 wt.% S) obtained with a reduced coating time (160°C, 3\nh).\nFinally, we are actively exploring process-intensification routes that could deliver both speed and\nuniformity. For example, preliminary trials of plasma-assisted sulfur coating (completed within 15\nminutes) could also rapidly yield uniform and stable coatings (Figure R4), suggesting a promising\n17 / 45\nalternative for scale-up. Other scalable approaches such as vapor-assisted spray or controlled vacuum\nevaporation in a tightly sealed chamber might also be effective for further process intensification.\nFigure R4. TEM image of L6CO@S prepared by plasma-assisted sulfur deposition.\nComment:\n2. For spent cathode material, it includes electrolyte, CEI, as well as binders, which has a high content\nof fluorine. I wonder its reaction pathway during flash upcycling method, its potential environmental\neffects and how to solve them.\nResponse:\nThank you for the insightful question on the potential environmental impact of fluorine-containing\ncomponents in spent cathodes. Indeed, the major sources of fluorine in aged electrodes are the residual\nelectrolyte salt (LiPF ), the cathode–electrolyte interphase (CEI), and the PVDF binder.\n6\nin an argon-filled glovebox, after which the electrodes were thoroughly rinsed with dimethyl carbonate\n(DMC). This step not only prevents LiPF from contacting ambient moisture, but also enables the\n6\ncollected electrolyte solution to be directly processed and recycled using established solvent and salt\nrecovery procedures. As a result, LiPF is effectively removed at the disassembly stage, avoiding\n6\nuncontrolled hydrolysis of LiPF and HF release.\ncomponents undergo rapid thermal decomposition. To demonstrate this, we carried out a control\nexperiment in which commercial LCO was mixed with PVDF and Li O and processed under identical\ncomplete conversion to L6CO (Figure R5), indicating that PVDF provides reductive carbonaceous\nspecies rather than persisting as an impurity. Consistent with this, FTIR of SLCO displays characteristic\nabsorptions peaks corresponding to PVDF (900–1400 cm−1) , which almost completely disappear after\nFJH (Figure R6). These results support the conclusion that binder residues are fully decomposed and\n18 / 45\nnot retained in the final product.\nFigure R5. XRD patterns of products obtained from LiCoO when PVDF was used as the sole reductant\n2\nunder FJH (1200°C, 10 s).\nFigure R6. FTIR spectra of SLCO before and after FJH conversion to L6CO.\nBecause the FJH reaction is extremely brief, the fluorinated volatiles released are transient and\nconcentrated in the off-gas stream, which makes them amenable to engineering control. In an industrial\nimplementation, the exhaust stream can be directed through standard scrubbing systems, such as\nalkaline absorption towers, which are routinely employed to handle HF in the fluorochemical and\nsemiconductor industries. This ensures that fluorinated emissions are efficiently neutralized and do not\npose an environmental hazard. In addition, the overall fluorine content in the CEI and binder is relatively\nlow compared to the bulk electrode mass, meaning that the absolute quantity of fluorinated gases\ngenerated per kilogram of recycled cathode is limited. When combined with proper gas treatment, the\nenvironmental burden is therefore minimal.\nWe have included the FTIR data into the revised manuscript and added clarifications of the\n19 / 45\nenvironmental effect and the mitigation strategies (Page 6, lines 20–24, highlighted):\n“FTIR of SLCO displays characteristic absorptions peaks corresponding to PVDF (900–1400\ncm⁻¹), which almost completely disappear after FJH (Supplementary Fig. 13), confirming its full\ndecomposition. Owing to the brief reaction duration, the released fluorinated volatiles are transient and\nconcentrated in the off-gas stream, which makes them amenable to engineering control.”\nComment:\n3. Another problem related to spent cathode is the impurities, such as Al and Cu, which can negatively\naffect the electrochemical performance as demonstrated in the previous FJH work. I wonder the contents\nof these impurities and if these impurities can affect the performance of L6CO.\nResponse:\nThank you for raising this important point regarding metallic impurities. Residual aluminum and\ncopper from current collectors are a potential concern in the recycling of spent cathodes. In our work,\nwe employed a water electrolysis separation (WES) method to detach the active material from the Al\nfoil and Cu foil current collectors under mild electrochemical conditions. This process allows the foils\nto be removed intact (Figure R7), rather than fragmenting into the powder, thereby minimizing the risk\nof metallic contamination. Consistent with this, inductively coupled plasma (ICP) analysis (Figure R8)\nshows that the concentration of Al in the recovered SLCO powder is only 1.8 ppm, corresponding to a\nmolar fraction below 1%, while Cu impurities are nearly undetectable.\nMoreover, because L6CO is used only as a lithium-supplementing additive in our full cells, its\nloading in the cathode is limited to ~4 wt.% of the total electrode mass. Even if trace Al or Cu impurities\nwere present, the absolute amount introduced into the composite electrode would be negligible relative\nto the bulk material. This is corroborated by our electrochemical testing: cells containing L6CO@S\nadditives display stable cycling performance without signatures of parasitic reactions, confirming that\nany residual Al or Cu does not measurably impact cell operation.\n20 / 45\nFigure R7. a, Photograph of the WES setup. b, Optical image of the SLCO electrode before WES. c,\nRecovered Al foil and separated SLCO cathode powder after WES.\nFigure R8. ICP results showing the concentrations of Al and Cu impurities in SLCO.\nTaken together, these results demonstrate that our separation procedure effectively suppresses\nAl/Cu contamination, and that the trace levels present after upcycling are too low to compromise the\nelectrochemical function of L6CO in practical applications. We have included the ICP results in the\nrevised manuscript to clarify these points (Page 4, lines 15–17, highlighted):\n“ICP analysis further confirmed that the residual Al content in the SLCO powder was only 1.8 mg\nL⁻¹, and no Cu impurity was detected (Supplementary Fig 2).”\nComment:\n21 / 45\n4. “Based on a comprehensive assessment of reaction kinetics and energy efficiency, 1200°C for 10\nseconds was identified as the optimal synthesis condition (Supplementary Fig. 7).” Can the authors\nexplain more details about reaction kinetics and energy efficiency, I cannot get the rationale.\nResponse:\nThank you for your insightful question regarding the choice of optimal synthesis condition. At\ntemperatures below 1000°C, the solid-state reaction proceeds too slowly to achieve full conversion\nwithin the ultrafast heating window. XRD patterns of products obtained at 800–1000°C exhibit\npronounced residual LiCoO peaks (Figure R9). When the reaction temperature is increased into the\n2\nrange of 1100–1200°C, the transformation kinetics improve significantly. At 1200°C, the 10 s pulse is\nsufficient to drive a complete rearrangement into the anti-fluorite structure. XRD results confirm phase\npurity, while SEM images in our manuscript confirm the complete removal of PVDF residues. However,\nraising the peak temperature above 1300°C does not further accelerate the kinetics but may instead lead\nto unnecessary energy input.\nWe also investigated the influence of heating duration at 1200°C. As shown in Figure R10, a 5 s\npulse still leaves detectable LiCoO , whereas complete conversion to Li CoO is achieved after 10 s.\n2 6 4\nExtending the dwell time to 20 s does not induce further phase changes, confirming that 10 s is sufficient\nfor full transformation.\nFigure R9. XRD patterns of the products obtained from FJH of SLCO at different temperatures for 10\nseconds.\n22 / 45\nFigure R10. XRD patterns of products obtained from FJH of SLCO at 1200°C for different reaction\ndurations.\nTaken together, 1200°C for 10 s represents the best compromise. We have included the above\nXRD data and clarified this rationale in the revised manuscript (page 5–6, lines 22–23 and lines 1–5\nhighlighted):\n“At 1200 °C, a 10 s pulse enables complete conversion into phase-pure Li₆CoO₄, whereas further\nincreases in temperature or dwell time do not improve the transformation and instead incur unnecessary\nenergy input (Supplementary Fig. 8). Thus, 1200°C for 10 seconds was identified as the optimal\nsynthesis condition, balancing transformation kinetics with energy efficiency. The corresponding\ntemperature–time profile during the FJH process (Supplementary Fig. 9) shows that the sample\ntemperature rises sharply to 1200°C, confirming the ultrafast heating characteristics of this technique.”\nComment:\n5. “The LFP+4% L6CO@S||Gr cell delivered outstanding initial capacities of 205.2 mAh g−1 (charge)\nand 159.3 mAh g−1 (discharge)”. I wonder if there are >10% extra Li stored at the anode side, will the\nfirst charging cycle cause lithium metal deposition at the anode side, since the N/P ratio is 1.1?\nResponse:\nThank you for the critical comment. In our study, the N/P ratio was carefully tuned to accommodate\nthe sacrificial lithium released by L6CO@S. Specifically, blank LFP||Gr cells were assembled with an\nN/P ratio of 1.1, while cells incorporating the lithium supplement were intentionally designed with an\nN/P ratio of 1.2 to further mitigate the risk of Li deposition. It is important to note that our N/P ratio is\ncalculated on the basis of the reversible capacity (Figure R11). The graphite anodes employed in this\nwork deliver an initial discharge capacity of ~380 mAh g−1 and an initial charge capacity of ~340 mAh\ng−1; if the irreversible capacity is taken into account, the effective N/P ratio becomes 1.35. When lithium\nreleased from both the LFP cathode and the L6CO@S additive is considered, the N/P ratio remains at\n~1.06. Only when the total first-charge capacity of the cathode exceeds 213.6 mAh g−1 would the\ngraphite approach its reversible storage limit. Consistent with this analysis, the first-charge voltage\n23 / 45\nprofiles exhibit no plateau indicative of metallic Li formation.\nFigure R11. Initial voltage–capacity curves of LFP, LFP+4% L6CO@S, and graphite electrodes in half-\ncell configuration, used to clarify the calculation of N/P ratio.\nWe have clarified the calculation of N/P ratio in the revised manuscript (page 18, lines 10–11,\nhighlighted):\n“The N/P ratio was calculated on the basis of the reversible capacities of the electrodes.”\nComment:\n6. S is an insulator and substantial amount of conductive carbon is required in Li-S batteries. Here, S\ncoating is used, and I wonder why it does not harm the specific capacity and rate performance?\nResponse:\nThank you for your thoughtful question. While bulk sulfur is indeed insulating, the coating\nemployed in this study is thin and conformal, optimized to passivate the surface and suppress lattice-\noxygen release without disrupting charge transport. In composite electrodes, electrons percolate\nprimarily through the carbon network and particle–particle contacts, a nanometer-scale sulfur film does\nnot sever these pathways and therefore does not fundamentally block electron percolation.\nAs shown in Figure R12, the first-charge curves of pristine L6CO and L6CO@S (10 wt.% S)\ndisplay comparable polarization, indicating that the modest capacity decrease in the coated sample is\nnot primarily conductivity-driven. Instead, the difference mainly arises from altered reaction pathways\nintroduced by the sulfur coating, where a portion of active lithium reacts to form stable lithium sulfate,\nthereby slightly reducing the accessible capacity. Prior studies have also shown that sulfur-containing\ninterphases at nanometer thicknesses can stabilize cathodes or anodes without compromising rate\nperformance1-3, supporting the view that when the sulfur-based layer is thin and uniform, it effectively\npassivates surface reactions while leaving both electron percolation and Li⁺ transport intact.\n24 / 45\nFigure R12. Initial charge curves of L6CO and L6CO@S at 0.05 C.\nTo further evaluate the kinetic influence of the sulfur coating, we performed electrochemical\nimpedance spectroscopy (EIS) and galvanostatic intermittent titration technique (GITT) measurements.\nEIS results (Figure R13) show that the charge-transfer resistance of L6CO@S is comparable to that of\npristine L6CO, confirming that the sulfur coating does not significantly increase electrode impedance.\nMoreover, since the additive accounts for only 4 wt.% of the total cathode (with sulfur contributing\nmerely 0.4 wt.%), its influence on the overall electrode resistance is negligible. Meanwhile, GITT-\nderived Li⁺ diffusion coefficients (Figure R14) for both pristine L6CO and L6CO@S remain in the\nsame order of magnitude across the practical state-of-charge range, indicating that lithium-ion transport\nis essentially preserved at the 10 wt.% sulfur loading used in this study.\nFigure R13. a, Electrochemical impedance spectroscopy (EIS) spectra of L6CO||Li and L6CO@S||Li\nhalf cells. b, EIS spectra of LFP||Li, LFP+4 % L6CO||Li, and LFP+4 % L6CO@S||Li half cells.\n25 / 45\nFigure R14. a,b, Voltage profiles of GITT measurements for L6CO and L6CO@S. c, Corresponding\nLi⁺ diffusion coefficients derived from the GITT data.\nThese results show that a thin and conformal sulfur coating does not impair either electronic\npercolation or ionic transport. The GITT data have been added to the revised manuscript and\nSupplementary Information (Page 8, lines 4–7), along with experimental methods (Page 18, lines 17–\n19, highlighted):\n“Moreover, GITT measurements reveal that the Li⁺ diffusion coefficients (D ) of pristine L6CO\nLi\nand L6CO@S remain in the same order of magnitude before and after coating (Supplementary Fig. 26),\nindicating that lithium-ion transport is not significantly hindered by the sulfur layer.”\n“The GITT measurements were conducted with the titration step at 0.1C for 20 min and a\nrelaxation step of 1 h.”\nComment:\n7. I wonder by adding extra L6CO@S, what the increasing ratios of specific capacity in full pouch cells\nare?\nResponse:\nThank you for this practical question. To quantify the benefit, we compared LFP||Gr pouch cells\nwith and without the addition of L6CO@S. Under the same formation and cycling protocol, the cells\ncontaining L6CO@S delivered an initial discharge specific capacity about 12.6% higher than that of\n26 / 45\nthe baseline cells (Figure R15), and this improvement was sustained during extended cycling. After\n1400 cycles at 0.5C, the capacity enhancement increased to 38.7%.\nFigure R15. Capacity improvement ratios of pouch cells with and without L6CO@S over different\ncycle numbers.\nThis improvement originates from the sacrificial lithium released by L6CO@S, which\ncompensates for the irreversible loss associated with SEI formation on the graphite anode, maintains\nhigher active Li inventory, and stabilizes the electrode–electrolyte interface. Collectively, these effects\ntranslate into substantially higher energy density and extended cycle life without altering the cell design\nor increasing electrode loading. We have added this detailed comparison to the revised manuscript to\nhighlight the practical performance gains enabled by our approach (Page 14, lines 13–16, highlighted):\n“As shown in Fig. 4g, LFP||Gr pouch cell incorporating L6CO@S delivers an initial discharge\ncapacity that is 12.6% higher than that of control cell, and the improvement becomes more pronounced\nwith extended cycling. After 1400 cycles at 0.5C, the capacity enhancement reaches 38.7%, with an\noverall capacity retention of 91.4%.”\nComment:\n8. More details about the flash upcycling method should be provided, including reactant mass,\nreactor dimensions and so on. Since the heat and mass transfer are directly related to reaction condition,\nthe upscaling ability should be demonstrated.\nResponse:\nThank you for your suggestions. We have provided more experimental details for the flash upcycling\nprocess in the revised manuscript (Pages 17–18, highlighted).\na) Disassembly and pre-treatment of spent pouch cells:\n“Spent LiCoO (SLCO) batteries were first discharged to 1.0 V and then soaked in 1.0 M NaCl\n2\naqueous solution for 24 h to ensure complete discharge. Disassembly was performed inside an argon-\n27 / 45\nfilled glovebox. Recovered electrode foils were rinsed with dimethyl carbonate (DMC) to remove\nresidual electrolyte and lithium salts, then dried under inert atmosphere prior to further processing.”\nb) Water electrolysis-induced separation (WES) process:\n“The recovered cathodes were directly employed as working electrodes in a two-electrode\nelectrolysis cell, with a mixed-metal oxide (MMO) mesh as the counter electrode owing to its high\noxidation resistance. An aqueous solution composed of 0.5 M Na SO with 0.01 M NaOH was\n2 4\nemployed as electrolyte. The electrolysis was conducted with a controlled current density of 50 mA\ncm−2, and the process was terminated once the active material was completely detached. The separated\nmaterials were subsequently washed with deionized water, dried at 80°C for 8 h, and gently ground.\nThe obtained black mass consists of a mixture of SLCO, conductive carbon, and polyvinylidene fluoride\n(PVDF).”\nc) Flash Joule heating (FJH) upcycling and sulfur coating process:\n“To convert SLCO to Li CoO , 0.291 g of spent cathode powder was mixed with 0.235 g of Li O\n6 4 2\nand ground in an agate mortar for 15 min. The powder mixture was then pressed into 14 mm diameter\npellets, sandwiched between carbon paper, and processed in a Joule heating device (Ultrafast High-\nTemperature Furnace 2022 A, Particle Precision instruments Co. Ltd, jilin Province, China). The\ntemperature was raised to 1200 °C within milliseconds and held for 10 s to yield Li CoO . For sulfur\n6 4\ncoating, the as-prepared Li CoO powder was mixed with elemental sulfur at a mass ratio of\n6 4\nLi CoO :S=10:1, hand milled for 15 min, sealed in a hydrothermal reactor under Ar, and heated in a\n6 4\nvacuum oven at 160°C for 6 h to produce L6CO@S.”\nRegarding scalability, flash Joule heating can be compatible with continuous, flow-type shock\nheating processes (Figure R16). In such configurations, spent cathode powder can be loaded onto a\nmoving conveyor belt, with current applied directly through the powder layer to induce rapid and\nuniform Joule heating. This approach has already been demonstrated in related shock-heating systems4-\n6, where conveyor-belt–assisted operation enables continuous, large-scale, and rapid high-temperature\ntreatment of powders. By adapting this concept to the present FJH process, scale-up to industrially\nrelevant throughputs is technically feasible.\nFigure R16. Schematic illustration of the industrial-scale FJH process for SLCO upcycling to L6CO.\n28 / 45\nComment:\n9. The LCA only considers the formation of L6CO, but the experiment parts focus more on L6CO part,\nI wonder if the whole procedure can be considered for the upcycling method. And how is the price of\nL6CO estimated, since it should be highly performance and quality-determined?\nResponse:\nThank you for your valuable suggestion. In our original study, the life-cycle boundary was defined\nfrom spent LCO electrode preparation through the flash-Joule-heating conversion to L6CO, which\nrepresents the dominant energy and value-adding step. The subsequent sulfur-coating step was excluded\nfrom the LCA boundary because its energy and environmental impact are negligible compared to the\nultrafast FJH process. Specifically, the coating is conducted in a low-temperature vacuum oven\n(<200 °C), which can process multiple samples simultaneously with a short dwell time (~6 h) and no\ngaseous emissions. The energy consumption of this mild heating process is orders of magnitude lower\nthan that of FJH, and thus does not materially affect the overall assessment. Moreover, our LCA\nframework was designed to benchmark the core upcycling pathway, ensuring consistent system\nboundaries across different scenarios.\nFor the application stage, this step simply involves blending a small mass fraction of L6CO@S\ninto standard cathode slurry. It does not modify the electrode manufacturing line, introduces no\nadditional chemical processing, and contributes negligibly to cost or environmental burden. Therefore,\nit was treated as outside the recycling LCA boundary.\nFor the economic comparison, L6CO is not yet a commercial product and therefore lacks a market\nprice. We therefore selected Li FeO as the comparator because it is one of the most representative,\n5 4\nproduction-ready lithium-rich sacrificial additives reported to date and is available at scale7. Although\nlithium oxalate (Li C O ) has been considered as a lower-cost lithium source, its high decomposition\n2 2 4\npotential (~4.5 V) typically requires the use of expensive catalytic additives to promote oxidation,\nmaking it impractical for typical cell formation workflows. By contrast, both L6CO and Li FeO exhibit\n5 4\nsuitable decomposition voltages, high irreversible capacities, and comparable anti-fluorite–type\ncrystal chemistry. In our electrochemical tests, L6CO delivers equal or slightly superior performance to\nLi FeO . Accordingly, using Li FeO provides a conservative, functionally relevant benchmark for\n5 4 5 4\neconomic assessment. We have clarified this rationale in the revised manuscript (Page 16, lines 9–16,\nhighlighted):\n“Notably, Li CoO can be synthesized at an estimated cost of ~$16,000 per ton. Since Li CoO is\n6 4 6 4\nnot yet commercially available, we benchmarked its price against Li FeO (~$69,500 per ton), one of\n5 4\nthe most representative lithium-rich sacrificial additives already produced at scale. Both Li CoO and\n6 4\nLi FeO share an anti-fluorite structure and exhibit suitable decomposition voltages with high\n5 4\nirreversible capacities, with Li CoO delivering slightly superior electrochemical performance. This\n6 4\n29 / 45\npronounced cost–performance advantage highlights the commercial viability and differentiated value\nof Li CoO within next-generation battery manufacturing.”\n6 4\nComment:\n10. Recent spent batteries have diverse chemistry, like NMC, NCA or their mixture within spent\ncathodes, how can the flash upcycling method be used in these spent cathodes?\nResponse:\nWe appreciate your insightful question regarding the adaptability of our FJH route to cathodes\nbeyond LiCoO . To probe this possibility, we subjected NCM523 cathode to identical or slightly\n2\nmodified FJH conditions. Unlike the straightforward conversion observed for SLCO, the products from\nNCM523 did not crystallize into a single, well-defined lithium-rich phase. XRD patterns show neither\npure anti-fluorite phases nor a defined mixture of Li NiO , Li CoO , and Li MnO (Figure R17). Instead,\n2 2 6 4 6 4\nthe reflections are dominated by metallic Ni, residual lithium oxide and NCM523, partially lithiated\nmixed oxides, and spinel-like features, suggesting that the present current pulse parameters are\ninsufficient to fully reorganize the more complex multicomponent lattice. Prolonging the FJH duration\ndid not induce further phase evolution. Electrochemical tests of the obtained materials yielded charging\ncapacities of 238–286 mAh g−1 and a discharging capacities of 17–26 mAh g−1 (Figure R18),\nindicating partial lithium-supplementing ability but falling short of an ideal lithium supply additive.\nFigure R17. XRD patterns of products obtained by FJH treatment of NCM523 mixed with\nLi O and carbon black at 1200°C for various durations. The molar ratio of the reactants is\n2\n30 / 45\nFigure R18. Initial charge–discharge curves of the products obtained by FJH treatment of NCM523\nmixed with Li O and carbon black at 1200 °C for different durations.\nheating. The presence of Ni, Mn and Co complicates rapid lattice reorganization and tends to promote\nphase separation rather than the clean formation of a single Li-rich compound. Nevertheless, we view\nthis as an important and promising research direction. Prior studies have shown that elements such as\nZn and Fe can co-occupy tetrahedral sites within anti-fluorite lattices, producing stable multimetal\nlithium-rich phases with attractive lithium-supplementing properties8,9. Drawing on this precedent, we\nbelieve that with optimized stoichiometry, tailored pulse parameters, and potentially staged or stepwise\nheating protocols, the FJH approach can be extended to NMC, NCA, or even mixed-cathode scrap.\nThis would be an important direction for our follow-up work.\nReferences:\n1. Zhang, K. et al. Sulfuration of Li‐Rich Mn‐Based Cathode Materials for Multianionic Redox and Stabilized\nCoordination Environment. Adv. Mater. 34, 2109564 (2022).\n2. Tu, S. et al. Fast-charging capability of graphite-based lithium-ion batteries enabled by Li P-based crystalline\n3\nsolid–electrolyte interphase. Nat. Energy 8, 1365–1374 (2023).\n3. Tian, M. et al. Designer Lithium Reservoirs for Ultra‐Long Life Lithium Batteries for Grid Storage. Adv. Mater.\n36, 2400707 (2024).\n4. Zhang, H. et al. Kinetics Dominated, Interface Targeted Rapid Heating for Battery Material Rejuvenation. Adv.\nEnergy Mater. 15, 2404838 (2024).\n5. Zhang, H. et al. Transient and dry recycling of battery materials with negligible carbon footprint and roll-to-\nroll scalability. Energy. Environ. Sci. 16, 2561-2571 (2023).\n6. Shi, W. et al. Roll-to-roll synthesis of multielement heterostructured catalysts. Nature Synthesis 4, 836-847\n(2025).\n31 / 45\n7. Liu, C. et al. Air-stable Li FeO additive enabled by carbon coating for energy-dense lithium-ion batteries. Nat.\n5 4\nCommun. 16 (2025).\n8. Wen, N. et al. Li Zn Co O with high oxygen oxidation reaction activity as a cathode pre-lithiation material\n6 0.5 0.5 4\nfor lithium-ion batteries. Nano Energy 142 (2025).\n9. Lee, W. et al. Modulating Anion Redox Reactions and Structural Evolution Through Fe‐Substitution in Li CoO\n6 4\nHyper‐Lithiated Sacrificial Cathodes. Adv. Energy Mater. 13, 2302316 (2023).\n32 / 45\nResponse to Reviewer #3\nComment:\nIn this study, the authors propose a novel upcycling method for converting spent LiCoO into Li CoO\n2 6 4\nvia flash Joule heating, achieving complete lithium and cobalt recovery within 10 seconds. The\nsynthesized Li CoO is further sulfur-coated to mitigate air sensitivity and oxygen release during\n6 4\ndelithiation, ultimately serving as a high-capacity sacrificial additive to offset active lithium loss. The\nwork demonstrates promising electrochemical performance when integrated into LiFePO ||Graphite\n4\npouch cells, achieving 91.4% capacity retention over 1400 cycles. The manuscript is commendable for\nits well-integrated narrative that connects material design, mechanistic understanding, and\nelectrochemical validation. However, certain aspects require further clarification, particularly regarding\nmaterial synthesis, the redox mechanism of oxygen anions and sulfur, and battery cycling performance.\nOverall, this is a meaningful work and minor revisions are recommended:\nResponse:\nThank you very much for your thoughtful evaluation of our manuscript. We are delighted that you\nrecognize the novelty of our upcycling strategy. In response to your constructive comments, we have\nmade the following focused revisions:\na) We clarified the FJH conversion of SLCO to Li CoO , and briefly explained why CoO cannot\n6 4\nbe fully converted under identical conditions.\nb) The redox mechanism of sulfur was expanded with Capillary electrophoresis results and added\nreferences, providing evidence for sulfate formation and a cautious interpretation of transient\nintermediates.\nc) More details on the sulfur coating process (temperature, loading, amorphization mechanism)\nwere added, along with electrochemical evidence (charge profiles, rate tests, GITT) showing\nthat a thin conformal sulfur layer stabilizes the interface without impeding kinetics.\nd) The long-term effect of Li CoO @S on cycling was clarified.\n6 4\ne) The figure citation has been corrected (former Fig. 3i now Fig. 3f), and captions were updated\naccordingly.\nPlease see the following response for details. The cited text from the manuscript and\nSupplementary files is marked in blue in this response, while all revisions in the manuscript and\nSupplementary files are highlighted in yellow. We sincerely hope that these changes adequately address\nyour concerns and further strengthen the manuscript for publication.\nComment:\n33 / 45\n1. The waste LiCoO particles exhibit Co O on the surface, with the bulk material remaining LiCoO .\n2 3 4 2\nHow is full conversion of LiCoO to Li CoO achieved via flash Joule heating, and how is stoichiometry\n2 6 4\nof LixCoyO accurately determined? Why can’t CoO be fully converted to Li CoO by Joule heating?\n2 6 4\nResponse:\nThank you for your insightful questions, which we address below from three perspectives:\n1) Conversion of Co O to Li CoO under FJH\n3 4 6 4\nSurface Co O does not hinder the complete transformation of LiCoO to Li CoO (L6CO) during\n3 4 2 6 4\nFJH. As part of this work, we previously explored the feasibility of upgrading spent LCO with different\ndegradation levels by simulating them through mixing fresh LCO with defined amounts of Co O ,\n3 4\nincluding a 100% Co O control as an extreme case without residual Li. The reaction pathway and\n3 4\nproduct XRD patterns are shown in Supplementary Fig. S17. Under the same FJH protocol, phase-pure\nL6CO was obtained in all cases, indicating that Co O can be fully converted to L6CO under these\n3 4\nconditions.\n2) Determination of Li Co O stoichiometry\nx y 2\nICP-OES was used to determine the stoichiometry of spent LCO. Since cobalt loss during cycling\nis typically minimal, we assumed y≈1 for SLCO and focused on determining the Li:Co ratio. Ratio\ndetermination is generally more accurate and reliable than absolute concentrations, and the Li:Co ratio\nalone is sufficient to calculate the precursor required for upcycling SLCO to L6CO.\nIn addition, powder XRD combined with Rietveld refinement was employed to quantify phase\nfractions. For example, refinement of one SLCO sample (Fig. XX) indicated ~9.3 wt.% Co O\n3 4\n(corresponding to a Li:Co ratio of ~0.889), while ICP-OES analysis of the same specimen gave a Li:Co\natomic ratio of 0.882. The close agreement between the structural (XRD) and compositional (ICP-OES)\nresults confirms the reliability of our stoichiometric determination and is consistent with established\npractice for analyzing spent electrodes1,2.\n3) Converting CoO to L6CO\nOur previous results showed that, unlike LiCoO feedstocks, CoO was not completely converted\n2\nto L6CO under identical FJH conditions (1200°C for 10 s, vacuum), and LiCoO impurities were\n2\ndetected. Similar oxidation behavior of CoO has also been reported by Prof. Ceder and co-workers3. In\ncontrast, conventional tube furnace heating under Ar (e.g., 700°C for 12 h) yielded phase-pure L6CO\n(Figure R1).\nWe attribute the incomplete conversion during FJH to residual oxygen: the base pressure\nachievable in our FJH chamber is approximately −0.1 MPa, corresponding to an absolute pressure of\n1.3 kPa (~98.7% evacuation), which cannot ensure a rigorously low-oxygen environment. Under such\nconditions, trace oxygen can interfere with the reaction and form LiCoO . Importantly, adding a small\n2\n34 / 45\namount of carbon black (~1 wt.%) effectively scavenged residual oxygen, after which phase-pure L6CO\nwas obtained (Figure R1). For clarity and to avoid overgeneralization, we have removed the earlier\nstatement implying that CoO is intrinsically more difficult to convert under FJH conditions.\nFigure R1. XRD patterns of Li CoO synthesized using CoO as the precursor. The products were\n6 4\nobtained by conventional solid-state reaction in an Ar-filled tube furnace (700 °C, 12 h), or by FJH with\nadditional 1 wt.% carbon black at 1200 °C for 10 s.\nComment:\n2. While the manuscript states that the modified Li CoO @S exhibits excellent rate performance, has\n6 4\nthis performance been evaluated when coupled with LiFePO ?\nperformance vary at different initial cycle currents?\nWhat is the observed decomposition rate under these conditions?\nResponse:\nThank you for your important question regarding the decomposition rate and efficiency of\nL6CO@S in LFP electrodes, which is more representative of practical use than testing the rate\nperformance of L6CO@S alone. To assess its oxidation behavior in LFP electrode, we evaluated the\nfirst-cycle performance of LFP+L6CO@S electrodes in half-cells and directly compared them with\nblank LFP electrodes. Current rates from 0.1C to 1C were applied, and the corresponding charge curves\nare shown in Figure R2. The additional capacity observed in the blended electrodes relative to blank\nLFP was attributed to the irreversible oxidation of the sacrificial L6CO@S. This excess capacity was\nnormalized by the L6CO@S mass to obtain an equivalent specific capacity, and the decomposition\nfraction was calculated by comparing this value with the specific capacity of L6CO@S electrodes tested\nindependently under the same rates.\n35 / 45\nFigure R2. a–d, Initial charge curves of LFP||Li and LFP+4% L6CO@S||Li half-cells at rates from 0.1–\n1C. e, Specific capacity calculated by the mass of L6CO@S at different rates and corresponding\ndecomposition ratios.\nAs shown in Figure R2e, the decomposition fraction of L6CO@S remains high across all tested\ncurrents, reaching ~96% even at 1C. These findings demonstrate that L6CO@S can be efficiently and\nnearly fully activated during the formation process, supporting its compatibility with fast formation\nprotocols. The corresponding data and analysis have been incorporated into the revised manuscript\n(Page 12, lines 24–28, highlighted):\n“To determine the real decomposition ratio of L6CO@S when coupled with LiFePO₄, additional\ntests were conducted at 0.1–1C in LFP+L6CO@S||Li half-cells (Supplementary Fig. 43), showing that\nthe additive is nearly fully oxidized even at 1C, confirming its compatibility with high-rate formation\nstep”\nComment:\n3. How is the sulfur coating uniformly applied to Li CoO ? Could the authors provide further details on\n6 4\nthe coating process? Given sulfur's poor conductivity, how does the sulfur coating impact the impedance\nof the material and its decomposition rate during cycling? Furthermore, how is the sulfur loading at\n160°C controlled to ensure an amorphous state?\nResponse:\nThank you for your insightful questions.\n1. Coating process and uniformity\n36 / 45\nThe sulfur coating was prepared by hand milling L6CO powders with microsized sulfur powder\n(1–2 μm) for 15 min to ensure uniform mixing, followed by heating the mixture at 160°C for 6 h in a\nsealed hydrothermal reactor under Ar. This approach is widely used for sulfur loading in Li–S cathode\nmaterials, where it enables efficient coating of sulfur onto the host surface4,5. Elemental sulfur melts at\n~115°C, and near 159°C it undergoes the λ-transition, during which S rings open to form long-chain\n8\npolysulfides. At this stage, the molten sulfur exhibits minimal viscosity and excellent flowability,\nallowing it to readily wet the L6CO particles and form a uniform surface layer. TEM images (Figure 2b\nand Supplementary Figure 21) of the coated particles confirm the presence of a continuous and\nhomogeneous sulfur coating on L6CO.\n2. Crystallinity of the sulfur layer\nUpon cooling from the λ-transition region, the long-chain polymeric sulfur species hinder long-\nrange ordering and suppress the crystallization of orthorhombic sulfur. As a result, the deposited sulfur\nlayer is predominantly amorphous or weakly crystalline rather than forming well-ordered crystalline\nsulfur. This interpretation is consistent with our XRD results, which show no distinct diffraction peaks\ncorresponding to crystalline sulfur. Similar amorphous or poorly crystalline sulfur phases have also\nbeen reported in many Li–S cathode materials prepared under comparable thermal conditions.\n3. Impact on impedance and decomposition rate\nWe acknowledge the concern that an insulating sulfur layer might hinder electrode kinetics.\nHowever, impedance measurements of L6CO@S (Figure R3) show that its impedance is comparable\nto that of pristine L6CO, indicating that the thin and conformal sulfur coating does not significantly\nalter the electrode resistance. Moreover, since the additive loading is small (4 wt.% of the total cathode,\nwith sulfur comprising merely 0.4 wt.%), the coating has negligible influence on the overall\nelectrochemical performance of the electrodes (Figure R3).\nRegarding decomposition behavior, the sulfur layer slightly affects the initial oxidation capacity.\nAs shown in Figure R4, the first-charge profiles of L6CO and L6CO@S exhibit comparable polarization.\nThe lower specific capacity of L6CO@S relative to pristine L6CO is primarily attributed to differences\nin the reaction pathway rather than conductivity limitations. This conclusion is supported by GITT\nmeasurements (Figure R5), which show that the Li⁺ diffusion coefficients of L6CO and L6CO@S\nremain in the same order of magnitude across the relevant state-of-charge range. In addition, rate-\nperformance tests further confirm that the sulfur coating does not impair decomposition kinetics. As\nshown in Figure R6, the charge capacities of pristine and coated electrodes decrease with increasing\ncurrent at a comparable rate, indicating that polarization between 0.1–0.5C is not aggravated by the\nsulfur layer.\n37 / 45\nFigure R3. a, Electrochemical impedance spectroscopy (EIS) spectra of L6CO||Li and L6CO@S||Li\nhalf cells. b, EIS spectra of LFP||Li, LFP+4 % L6CO||Li, and LFP+4 % L6CO@S||Li half cells.\nFigure R4. Initial charge curves of L6CO and L6CO@S at 0.05 C.\nFigure R5. a,b, Voltage profiles of GITT measurements for L6CO and L6CO@S. c, Corresponding\nLi⁺ diffusion coefficients derived from the GITT data.\n38 / 45\nFigure R6. Comparison of rate performance between L6CO and L6CO@S.\nWe have provided additional details on the coating procedure (Page 17, lines 39–41) and included the\nGITT data in the revised manuscript (Page 8, lines 4–7), along with experimental methods (Page 18,\nlines 17–19, highlighted):\n“For sulfur coating, the as-prepared Li CoO powder was mixed with elemental sulfur at a mass\n6 4\nratio of Li CoO :S = 10:1, hand milled for 15 min, sealed in a hydrothermal reactor under Ar, and heated\n6 4\nin a vacuum oven at 160 °C for 6 h to produce L6CO@S.”\n“Moreover, GITT measurements reveal that the Li⁺ diffusion coefficients (D ) of pristine L6CO\nLi\nand L6CO@S remain in the same order of magnitude before and after coating (Supplementary Fig. 26),\nindicating that lithium-ion transport is not significantly hindered by the sulfur layer.”\n“The GITT measurements were conducted with the titration step at 0.1C for 20 min and a\nrelaxation step of 1 h.”\nComment:\n4. At 40% and 60% state of charge (SOC), the S K-edge XANES spectra show peaks for sulfur and\nLi SO , along with an additional smaller peak. What compound is represented by this smaller peak?\n2 4\nAre there any intermediate phases that could influence the battery system? Moreover, aside from soft\nX-ray analysis, are there other methods used to confirm the formation of Li SO and provide direct\n2 4\nevidence of the reaction between sulfur and reactive oxygen to generate [S ]+ and [S ]+?\n8 5\nResponse:\nThank you for this insightful question. The additional smaller feature in the S K-edge XANES\nspectra at intermediate states of charge (≈40–60% SOC) is most reasonably attributed to a partially\noxidized sulfur species rather than the fully oxidized +6 state. In particular, sulfur in the +4 oxidation\nstate (e.g., sulfite, SO 2−) has been reported to exhibit similar spectral signatures6. Its position between\n3\nelemental sulfur and sulfate, combined with its appearance only at intermediate SOC and disappearance\n39 / 45\nat higher SOC, is consistent with the stepwise oxidation of amorphous sulfur fragments that ultimately\nyield stable Li SO as charging proceeds. Our XPS measurements further support this interpretation.\n2 4\nNotably, the +4 intermediate (Li SO ) is considered a stable CEI component7,8; it is poorly soluble and\n2 3\nhas not been reported to cause harmful side reactions. Moreover, in our system its concentration is low,\nand no adverse electrochemical effects were observed.\nTo corroborate the formation of Li SO beyond soft X-ray analysis, we also performed X-ray\n2 4\nphotoelectron spectroscopy (XPS) and capillary electrophoresis (CE) measurements. XPS spectra\nclearly reveal the valence evolution of sulfur during charging, consistent with the progressive oxidation\npathway inferred from XANES. CE, owing to its high sensitivity to SO 2− (detection limit ~0.1 ppm in\n4\nour UV-based setup), enables accurate quantification of sulfate species. After disassembling the post-\ncharging cells, the L6CO@S electrodes were thoroughly washed, the supernatant was filtered, and the\nSO 2− concentration was measured. The detected sulfate concentration was approximately 39 ppm, as\n4\nshown in Figure R7 (originally Supplementary Fig. 33 of our prior submission). This value is in close\nagreement with the theoretical concentration expected for complete sulfur oxidation, thereby providing\nindependent confirmation of Li SO formation and supporting the high completeness of the reaction.\n2 4\nFigure R7. Quantification of SO 2− generated in charged L6CO@S electrodes by capillary\n4\nelectrophoresis. The measured value corresponds to 97.6% of the theoretical amount.\nWith respect to direct detection of radical-cation intermediates such as [S ]+ and [S ]+,\n8 5\nwe\nacknowledge the limitations of our current characterization. These species are highly reactive and short-\nlived, and their unequivocal observation generally requires advanced operando techniques (e.g., EPR\nor high-resolution MS)9, which are beyond the scope of this study. Nevertheless, the transient formation\nof sulfur radical cations is theoretically plausible. During charging, reactive oxygen species (such as\nsuperoxide or singlet oxygen) can be generated and may attack S–S bonds, leading to S ring opening\n8\nand the formation of chain-like sulfur intermediates (e.g., [S ]+, [S ]+). Such intermediates have been\n8 5\nreported in the literature under oxidative conditions. Accordingly, we have revised the manuscript text\nto present this point more cautiously and cite additional supporting studies.\nRevisions added to the manuscript (Page 11, lines 15–23 and Page 12, lines 7–9, highlighted):\n40 / 45\n“Between ~40% and 60% SOC, a intermediate feature appears between the positions of elemental\nsulfur and Li SO , consistent with the partially oxidized sulfur species (e.g., sulfites), which disappear\n2 4\nat higher SOC. Beyond 60% SOC, the edge shifts toward higher energy, with the emergence of features\nat 2478.6 eV and 2483.4 eV that are characteristic of sulfate species.”\n“Capillary electrophoresis was further employed to quantify SO 2− in fully charged L6CO@S\n4\ncathodes. The measured concentration (~39 ppm) closely matched the theoretical value (Supplementary\nFig. 33), providing independent confirmation of complete sulfur oxidation.”\n“Upon charging, S molecules may be attacked by reactive oxygen species, leading to S–S bond\n8\ncleavage and the formation of generating reactive chain-like sulfur intermediates (e.g., [S ]•+, [S ]•+,\n8 5\netc.).”\nComment:\n5. The manuscript suggests that the introduction of the prelithiation agent enhances both the discharge\ncapacity and cycling stability of the LiFePO ||Graphite full cell. However, the Li CoO prelithiation\n4 6 4\nagent primarily affects the first cycle and is unlikely to have a significant impact on subsequent cycles.\nPlease elaborate on the mechanism by which the prelithiation agent enhances cycling performance over\nsubsequent cycles. Additionally, since LiFePO ||Graphite cells are mature commercial systems with\n4\ngood cycling stability even without prelithiation, the significant capacity degradation observed in\nFigures 4b, 4g, and 4h requires further explanation.\nResponse:\nThank you for this perceptive question. We first clarify that LiFePO ||Graphite is indeed a mature\n4\ncommercial system with excellent cycling stability at the industrial level. However, laboratory-\nassembled cells generally cannot match the stability of commercial products, where long-term\nperformance is ensured by optimized assembly protocols, advanced electrode engineering, and\nproprietary electrolyte additives that are not easily reproduced under laboratory conditions. As is\ncommon in academic studies, our evaluation emphasizes internal comparisons under controlled\nconditions, where the only variable is the presence or absence of the prelithiation agent. This approach\nprovides a meaningful assessment of the effect of L6CO@S.\nRegarding the mechanism, it is important to emphasize that the loss of cyclable lithium is not\nconfined to the first cycle. The SEI on graphite is dynamic and continues to reform and thicken during\ncycling, steadily consuming lithium and driving gradual capacity fade. As shown in Figure R8, LFP\nhalf-cells display excellent stability, while graphite half-cells exhibit progressive capacity decay,\nprimarily due to ongoing electrolyte reduction and SEI reformation on newly exposed graphite surfaces.\nThis leads to small but persistent Coulombic inefficiencies that accumulate into measurable capacity\nloss over long-term cycling. For instance, Figure R9a shows that even with an average Coulombic\n41 / 45\nefficiency of 99.95%, the continuous active lithium loss (cALL) after 326 cycles is comparable in\nmagnitude to the initial active lithium loss (iALL).\nFigure R8. Cycling performance of LFP||Li and Gr||Li half-cells at 0.5C.\nOur strategy addresses this by introducing not only sufficient lithium to offset iALL but also an\nadditional lithium reservoir stored in the graphite anode (Figure R9b). This reserve is gradually\nmobilized to compensate for cALL during extended cycling, accounting for the slower capacity-decay\nrate observed in cells containing L6CO@S. Figure R9c further shows that the capacity retention of our\ncells closely follows decay curves predicted from measured Coulombic efficiencies, confirming that\ntheir degradation behavior falls within the expected range. In summary, the improvement in both\ndischarge capacity and cycle stability of cells with L6CO@S arises from two complementary effects:\n(i) the surplus lithium inventory provides a long-term buffer against continuous SEI-related\nconsumption, and (ii) the sulfur coating suppresses oxygen-driven side reactions that would otherwise\ndestabilize the anode SEI and accelerate fade.\nFinally, by benchmarking our results against representative LFP||Graphite cycling data reported in\nliteratures (Figure R10), we find that the capacity retention of our control cells is already superior to\nthat of most previously reported laboratory cells. This confirms that the improvements observed with\nL6CO@S are not due to an unduly weak baseline; rather, they represent a genuine enhancement built\nupon an already high level of cycling stability achievable under laboratory conditions.\n42 / 45\nFigure R9. a, Relationship between continuous active lithium loss (cALL) and cycle number under\ndifferent average Coulombic efficiencies (CE ). b, First-cycle voltage–capacity curves of LFP, LFP+4%\navg\nL6CO@S, and graphite electrodes. c, Predicted capacity retention curves at different CE compared\navg\nwith measured values of LFP||Gr and LFP+4% L6CO@S||Gr full cells.\nFigure R10. Comparison of the baseline LFP||Gr full-cell performance in this work with representative\nreports from the literature10-17.\nComment:\n6. Figure 3i is cited as comparing various prelithiation agents in terms of air stability and gas evolution,\n43 / 45\nbut this figure does not appear in the manuscript. Does the reference actually pertain to Figure 3f? Please\nclarify.\nResponse:\nThank you for the careful reading. We apologize for this oversight—Figure 3f was incorrectly cited\nas Figure 3i. This has been corrected in the revised manuscript (Page 12, line 16, highlighted).\nReferences:\n1. Zhang, S. et al. Adaptive Localized Ectopic Structure Enhances Regeneration of Spent Lithium‐Ion Battery\nCathodes. Adv. Mater. (2025).\n2. Wang, T. et al. Direct Recycling of Spent NCM Cathodes through Ionothermal Lithiation. Adv. Energy Mater.\n10 (2020).\n3. Jun, K. et al. Understanding the Irreversible Reaction Pathway of the Sacrificial Cathode Additive Li CoO .\n6 4\nAdv. Energy Mater. 13, 2301132 (2023).\n4. Hu, G. et al. A Sulfur‐Rich Copolymer@CNT Hybrid Cathode with Dual‐Confinement of Polysulfides for\nHigh‐Performance Lithium–Sulfur Batteries. Adv. Mater. 29 (2016).\n5. Li, H. et al. Kinetically‐Enhanced Gradient Modulator Layer Enables Wide‐Temperature Ultralong‐Life All‐\nSolid‐State Lithium–Sulfur Batteries. Adv. Energy Mater. 15 (2025).\n6. Zhang, K. et al. Sulfuration of Li‐Rich Mn‐Based Cathode Materials for Multianionic Redox and Stabilized\nCoordination Environment. Adv. Mater. 34, 2109564 (2022).\n7. Li, W. et al. Understanding and Design of Cathode–Electrolyte Interphase in High‐Voltage Lithium–Metal\nBatteries. Adv. Funct. Mater. 34 (2024).\n8. Chen, S. et al. Engineering durable interphases for high-voltage Li-ion batteries under thermal stress. Natl. Sci.\nRev. 12 (2025).\n9. Wang, Q. et al. Direct Observation of Sulfur Radicals as Reaction Media in Lithium Sulfur Batteries. J.\nElectrochem. Soc. 162, A474-A478 (2015).\n10. Chen, Y. et al. Converting Li‐Rich Layered Oxide Cathode into Non‐Shrinking Sacrificial Prelithiation Agent.\nAdv. Mater. (2025).\n11. Zhang, J. et al. Worse Interference of Fe3+ than Fe2+ on Degrading the Interphase and Performance of\nLiFePO ||Graphite Battery. Adv. Mater. (2025).\n4\n12. Sun, Y. et al. High-capacity battery cathode prelithiation to offset initial lithium loss. Nat. Energy 1, 15008\n(2016).\n13. Zhang, C. et al. Air-stable lithium-sandwiched current collector for non-destructive, thermally safe, and\nsustained supplementary lithiation. Energy. Environ. Sci. (2024).\n14. Chen, X. et al. Sustainable Prelithiation Strategy: Enhancing Energy Density and Lifespan with Ultrathin Li‐\nMg‐Al Alloy Foil. Adv. Energy Mater. (2024).\n15. Li, Z. et al. Electrolyte Design Enables Rechargeable LiFePO /Graphite Batteries from −80 °C to 80 °C.\n16. Zhang, Y. et al. Innovative Anode Design to Enhance Both Volumetric and Gravimetric Energy Densities of\n44 / 45\nLiCoO ||Graphite Pouch Cells. Adv. Energy Mater. 15 (2024).\n2\n17. Nie, Q. et al. Ketomalonate‐Based Lithium Replenishment Reagents for Lithium‐Ion Batteries With Stable\nElectrode Structural. Adv. Funct. Mater. (2025).\n45 / 45"}, {"round": 2, "reviewer_comments": {"reviewer_1": "The author has answered my questions, and I have no further questions. The article may be considered for acceptance.", "reviewer_2": "My concerns have been properly addressed by the authors, and I agree with the acceptance of this manuscript.", "reviewer_3": "The authors have thoroughly addressed all the concerns raised by the reviewer in the revised manuscript, and it is now\nrecommended for publication."}, "author_response": ""}]

Decision letter