Planar lithium deposition/dissolution enabling practical 500 Wh kg–1 anode-free pouch cells

Planar lithium deposition/dissolution enabling practical 500 Wh kg–1 anode-free pouch cells | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Physical Sciences - Article Planar lithium deposition/dissolution enabling practical 500 Wh kg–1 anode-free pouch cells Jianhui Wang, Lei Liu, Yuxuan Xiang, Xingyu Lu This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-5047161/v1 This work is licensed under a CC BY 4.0 License Status: Published Journal Publication published 17 Mar, 2026 Read the published version in Nature → Version 1 posted You are reading this latest preprint version Abstract Anode-free lithium metal batteries (AFLMBs), characterized by the absence of anode active materials during manufacturing, offer great potential for high-energy-density, low-cost energy storage. However, AFLMBs face a long-standing challenge of short lifespan due to the harsh conditions of lacking excess lithium and an anode host. This issue is associated with uneven lithium deposition/dissolution, rooted in the micro-inhomogeneity and fragility of solid electrolyte interphase (SEI) on the lithium metal surface. Here, we present a practical 500 Wh kg–1-level AFLMB design with enhanced lifespan, achieved using an electrolyte of 1.6 M lithium difluoro(oxalate)borate in N,N-Dimethyltrifluoroacetamide solvent. The electrolyte-derived B-F-based polymer-rich SEI exhibits sub-nanometer homogeneity, high flexibility, and fast Li-ion conductivity, which spontaneously evolves a self-adaptive mesh-film structure that ensures uniform ion flux and large-volume-change accommodation, thereby realizing reversible planar lithium-orientated deposition/dissolution of 5.6 mAh cm–2. Consequently, a 2.7 Ah AFLMB (508 Wh kg–1, 1668 Wh L–1) without any host-material coating demonstrates stable cycling for 100 cycles at 100% depth of discharge (DoD) and 250 cycles at 80% DoD, with 80% capacity retention and an unprecedentedly high-power output of 2650 W kg–1 at 96 Wh kg–1. Our findings address the inherent structural instability of host-free electrodes, advancing the practical implementation of AFLMBs. Physical sciences/Energy science and technology/Energy storage/Batteries Physical sciences/Materials science/Materials for energy and catalysis/Batteries Figures Figure 1 Figure 2 Figure 3 Figure 4 Introduction The quest for high-energy-density and low-cost rechargeable batteries has been a long-standing goal. After more than 30 years of commercial development, the energy density and cost of traditional lithium-ion batteries (LIBs) have nearly reached their theoretical limits 1-4 . The use of silicon (Si) to partially replace the graphite anode material in silicon-carbon (Si/C) composites can exceed the theoretical energy density limit of 280 Wh kg –1 for traditional LIBs. However, Si inherently suffers from poor electronic conductivity, low initial Coulombic efficiency, large volumetric expansion, and severe pulverization during cycling. These issues necessitate complex material engineering processes for practical application of high-specific-capacity Si/C anodes 5-8 . Consequently, the production cost of 360 Wh kg –1 Si-based LIBs is nearly double that of traditional LIBs, and mass production has yet to be realized. The lithium metal batteries (LMBs), which utilize lithium metal as the anode and high-nickel nickel-cobalt-manganese (NCM) as the cathode, offer an energy density exceeding 400 Wh kg –1 9,10 . However, the high chemical reactivity of metallic lithium and the challenges in processing ultra-thin lithium foil pose significant obstacles to the large-scale application of LMBs 11,12 , particularly in terms of production costs, safety, and product consistency. Moreover, both Si-based LIBs and LMBs are likely to experience a more pronounced increase in production costs as energy density rises (see Figure 1a ). In contrast, anode-free lithium metal batteries (AFLMBs), which do not use anode active materials during manufacturing, simplify battery configuration and reduce raw material usage 13-15 . This battery design offers the highest possible gravimetric and volumetric energy densities while reducing the production cost per unit of energy to a level even lower than that of traditional LIBs. Moreover, the cost of production does not significantly increase with higher energy density, positioning AFLMBs as an ideal approach for mass-producing high-energy-density batteries ( Figure 1a ). However, the lack of an anode host and excess lithium resource subjects AFLMBs to extremely harsh operating conditions, resulting in a shortened lifespan, especially under high-energy-density conditions 14,16 . To date, most AFLMBs have demonstrated limited cycling life (50-300 cycles) at relatively low energy densities (50-350 Wh kg –1 ) 17-25 , rendering them less competitive compared to traditional LIBs (160-280 Wh kg –1 , >800 cycles). Several attempts to develop AFLMBs at the 400 Wh kg –1 level have all showed a short lifespan of less than 80 cycles 16,26-29 . Currently, there have been no reports of repeatedly rechargeable AFLMBs achieving energy densities exceeding 500 Wh kg –1 . Design of practical 500 Wh kg –1 pouch cells Maximizing the active electrode materials while minimizing inactive materials is essential for constructing high-energy-density lithium batteries. In an AFLMB with a given cathode material, the energy density is determined by the cathode areal capacity, the number of stack layers, and the electrolyte weight to cell capacity (E/C), assuming that other technological parameters, such as current collectors, separator, tabs, and Al-plastic film, are established. The use of high areal capacity, multiple layers, and low E/C enhances battery energy density. However, excessive areal capacity and stack layers, or too low E/C, can introduce challenges in battery manufacturing and performances. Currently, commercial Ah-level LIBs with graphite anodes and NCM cathodes typically feature cell parameters of 5 ~ 20 cathode layers, 2.0 ~ 5.0 mAh cm –2 areal capacity, and 1.5 ~ 4.0 g Ah –1 E/C for optimal battery performances. Moreover, the current collectors for the anode (Cu foil, 4~12 μm thickness) and cathode (Al foil, 12~20 μm thickness), as well as the Al-plastic film (>88 μm thickness), are indispensable due to their chemical/mechanical stability and cost-effectiveness in Ah-level practical applications 30,31 . Therefore, to achieve the highest possible energy density under practical conditions, we selected 4.4 V NCM811 as the cathode material, owing to its highest specific energy density (216 mAh g –1 , 3.84 V average discharge voltage) among currently mature cathode materials, and minimized the weight of inactive materials by using 4 μm Cu foil, 12 μm Al foil, 12 μm polyethylene (PE) separator, 100 μm tabs, and 88 μm Al-plastic film. A mathematical model of pouch cells was established to correlate the cell energy density with key parameters of the cathode areal capacity, stack layer number, and E/C. The detailed calculations are provided in Table S2–S4 . As shown in Figure 1b , our calculated correlation matches well with the disclosed cell parameters of commercial LIBs (LGES, 268 Wh kg –1 graphite||NCM712 and Li-FUN, 200 Wh kg –1 graphite||NCM523), validating the accuracy of our model. According to the calculated AFLMB correlation, constructing a 500 Wh kg –1 AFLMB has a considerably higher requirement as compared to commercial LIBs, despite the absence of anode material. If E/C > 2.0 g Ah –1 , it requires more than 20 cathode layers with an areal capacity exceeding 6.8 mAh cm –2 , which is too high as compared to commercial NCM cathodes (2 ~ 5 mAh cm –2 ). If the cathode areal capacity is less than 4.0 mAh cm –2 , the E/C ratio would need to be below 1.3 g Ah –1 , which hardly ensures adequate wetting and rate capability regarding the porosity of electrodes and separator 32-34 . Furthermore, if the stack comprises fewer than 10 cathode layers, the weight percentage of Al-plastic film in the cell would be too high to reach the target. Considering all these factors, we ultimately selected cell parameters of 5.6 mAh cm –2 areal capacity, 10 stack layers, and 1.5 g Ah –1 E/C to construct a 500 Wh kg –1 -level AFLMB ( Figure 1 b, c ). These parameters are notably more stringent than those in previously reported AFLMBs (2 ~ 4 mAh cm –2 , 2 stack layers, and > 2.0 g Ah –1 ) 16-29 . The E/C ratio of 1.5 g Ah –1 approaches the lower limit for advanced LIBs 32-34 , imposing rigorous demands on the density, wettability, and stability of the electrolyte. Currently, only liquid electrolytes could meet these requirements. In this study, we selected an optimized electrolyte of 1.6 M lithium difluoro(oxalate)borate in N,N-Dimethyltrifluoroacetamide solvent (1.6 M LiDFOB/NDFA, named as BAFF electrolyte, illustrated in Figure 1d ) for the following reasons: (1) the electrolyte has a low density (1.36 g cm –3 ), viscosity (7.04 mPa s) and vapor pressure (1.15 kPa), as well as good wettability with cell materials at room temperature ( Figure S1a, Table S5 ); (2) it exhibits high ion conductivities across a wide temperature range of –40 to +60 ℃ (0.5 ~ 14.2 mS cm –1 , Figure S1b ); (3) it has a wide electrochemical stability window of ~5.0 V, enabling not only stable cycling of NCM811 to 4.5 V ( Figure S2-S3 ), but also highly reversible planar Li deposition/dissolution of 5.6 mAh cm –2 , thereby enhancing the longevity of 500 Wh kg –1 AFLMB (see discussed below); and (4) it produces almost no gas during cycling, overcoming the gassing issue generally faced by LiDFOB-based electrolytes ( Figure S3 ) 17,28 . Additionally, we tested AFLMBs using three advanced electrolytes reported in literature, including local high concentration electrolyte (LHCE), high concentration dual-salt electrolyte (HCDE), and all fluorinated electrolyte (AFE), for comparison. As depicted in Figure 1c-e , our designed AFLMB exhibits a total capacity of 2.7 Ah and total mass of 20.5 g, achieving a gravimetric energy density of 508 Wh kg –1 ( Table S6 ). Moreover, the absence of low-density graphite materials significantly reduces the cell thickness from 4.3 mm to 2.3 mm, resulting in an ultra-high volumetric energy density of 1668 Wh L –1 , which is more than double that of the LIB using the same cell components ( Figure 1c ). When excluding the Al-plastic film, the stack energy densities of our AFLMB reach 543 Wh kg –1 and 1806 Wh L –1 . Reversible planar lithium-orientated deposition/dissolution Although AFLMBs do not utilize active anode materials during manufacturing, a micrometer-scale layer of lithium metal is deposited on the anode current collector, whose stability during cycling dictates the battery lifespan. Figure 2a illustrates the initial five cycles of a Cu||NCM811 anode-free pouch cell using the BAFF electrolyte at a current density of 0.56 mA cm –2 . The Coulombic efficiency is 93.8% in the first cycle and rapidly increases to 99.8% in the third cycle, indicating highly reversible lithium deposition/dissolution under the high areal capacity of 5.6 mAh cm –2 . A reversible specific capacity of 216 mAh g –1 was obtained in the pouch cell, which is equal to that of Li||NCM811 half-cell, indicating a 100% capacity utilization. In sharp contrast, dendritic short-circuiting occurs in the conventional carbonate electrolyte (CME) even during the first charge under these harsh conditions (see Figure S4a ). To evaluate the quality of the deposited lithium on the bare Cu current collector, charged samples were collected after five cycles and analyzed by cryogenic focused ion beam scanning electron microscope (Cryo-FIB-SEM). The cross-sectional image ( Figure 2b ) shows a deposited lithium thickness of 28 μm, which closely matches the theoretic thickness of 27 μm for 5.6 mAh cm –2 of metallic lithium. The upper half of the deposited layer is composed of columnar grains approximately 8 μm in diameter, while the lower part exhibits no grain boundaries, demonstrating an integrated and compact structure. Through a substantial number of cross-sectional slice images, a 3D reconstruction of the deposited lithium with dimensions 10×28×5 μm reveals a porosity of only 1.6% for BAFF, which is remarkably superior to those obtained with reference electrolytes AFE, LHCE and HCDE, whose deposited lithium thicknesses are 37 μm, 35 μm, and 33 μm with corresponding porosities of 20.0%, 16.6%, and 11.1%, respectively (see Figure S5 ). Subsequently, the deposited lithium was subjected to 2D grazing-incidence X-ray diffraction measurement (2D GIXD). As shown in Figure 2c -upper, distinct flecks of the (110) crystal plane can be clearly observed for BAFF, demonstrating a Li (110) texture characteristic, which indicates that BAFF enables the orientated growth of metallic lithium on bare Cu substrate. This finding contrasts sharply with reference electrolytes (see Figure 2c -lower, S6 ) and prior studies 16,29 , wherein abundant lithium mosses typically grow on the Cu surface ( Figure S7a ) due to its lithiophobic and uneven nature. To induce lithium orientated growth on bare Cu substrate, it is necessary to firstly deposit lithium in a thickness >50 μm to overcome the adverse substrate effect 35 . Surprisingly, the use of BAFF directly grows lithium texture on bare Cu with minimal nucleation overpotential (8 mV, Figure 2a inset, S4c ) and lithium mosses (see bottom view in Figure S7a, S9a ), indicating that BAFF effectively mitigates the adverse impact of the Cu substrate. Further, scanning electron microscope (SEM) was employed to directly observe the morphology changes of deposited lithium at different states during cycling. In BAFF, initial charge to 3.6 V produced a uniform distribution of lithium nuclei with diameter of approximately 200 nm on bare Cu ( Figure 2d ), which grew into large grains with sizes of 5~8 μm upon charging to 4.4 V ( Figure 2e ). Interestingly, in the following discharge to 3.9 V, the dissolution of lithium resulted in a “hemispherical pit” on each grain ( Figure 2f ). Since the lithium metal surface is covered by a solid electrolyte interphase (SEI), the deep indentation morphology reflects that the SEI keeps firmly adhered to the shrinking lithium metal without fracturing, demonstrating exceptional flexibility and lithiophilicity of the SEI. After fully discharging to 1.0 V, all the deposited lithium metal dissolved, leaving behind a unique mesh film covering the Cu surface ( Figure 2g ). This is attributed to SEI aggregation caused by the full stripping of lithium columnar grains: the vertical contraction of SEI on the side surface of columnar grains forms the thick wall of mesh cell, while the horizontal downshift of SEI on the upper surface of columnar grains forms ultra-thin face of mesh cell, beneath which the Cu substrate can be observed via energy-dispersive spectroscopy (EDS) mapping ( Figure 2i, S7, S9 ). Additionally, the size of these mesh cells (5~8 μm) is consistent with that of lithium columnar grains, suggesting minimal change in lithium grain size during the dissolution process. Upon the second charge to 3.6 V, this SEI-aggregated mesh film re-bulged as lithium metal plating ( Figure 2h ), indicating a reversible and self-adaptive process. Notably, throughout the charge/discharge processes, all lithium grains exhibit nearly synchronous bulge/pit surface, demonstrating a typical planar (2D) deposition/dissolution characteristic ( Figure S8, S10 ). Meanwhile, the SEI remains firmly adhered to the lithium or Cu surface, effectively preventing electrolyte infiltration ( Figure 2g ). Both collectively ensure maximum structural stability during cycling, as strongly supported by EDS results ( Figure 2i ), which show a clear mesh-structure film on bare Cu for many cycles. By contrast, lithium deposition/dissolution in all the reference electrolytes shows very different behaviors ( Figure 2j, S7, S9 ). During discharge, lithium grains contract and the gaps between grains increase following a typical 3D dissolution process, leading to either porous SEI (e.g., LHCE and HCDE) or cracked SEI (e.g., AFE) ( Figure S7 ). This loose structure cannot prevent electrolyte infiltration, which inevitably creates numerous new sites for lithium nucleation and growth in subsequent charging, resulting in dynamically changing structure during cycling ( Figure S9 ). Consequently, the planar lithium deposition/dissolution unique to BAFF renders a stable mesh-film-structure SEI on lithium metal during cycling, effectively overcoming its inherent shortcoming of lacking a host. SEI characterization and mechanistic understanding To understand the mechanism behind the unique planar lithium deposition/dissolution, we conducted a detailed study of the composition, structure, and mechanical and ion-conductive properties of the SEI after five cycles. As shown in Figure 3a , cryogenic transmission electron microscopy (Cryo-TEM) revealed that the BAFF-derived SEI on the Li (110) grains is approximately 8 nm thick and in amorphous state (see fast Fourier transform (FFT) analysis in Figure 3a insert ). The nano-scale morphology and mechanical properties of the SEI were characterized by atomic force microscope (AFM). Figures 3b-e and S11 show that the SEI has an exceptionally smooth surface with an ultra-low roughness (Ra = 0.7 nm) and a relatively high Young’s modulus of 3.15 GPa with a very narrow distribution (Ra = 0.21 GPa). This character of ultra-high uniformity at sub-nanometer scale for the BAFF-derived SEI is in sharp contrast to the significant heterogeneity observed in the SEIs from the reference electrolytes (LHCE, HCDE, and AFE), which exhibited considerably rougher surface (Ra = 4.6, 3.3, and 9.2 nm, respectively) and more widely distributed Young’s modulus values (Ra = 0.75, 0.52 and 0.82 GPa, respectively). Regarding to the chemical composition of the BAFF-derived SEI, X-ray photoelectron spectroscopy (XPS) identified the primary elements (B, C, N, O and F) and some ingredients like organic B-F or C-F species, B-O species, and a small amount of LiF ( Figure S12 ). Due to the SEI’s susceptibility to air and its insolubility in various high-polarity aprotic solvents such as Dimethyl sulfoxide (DMSO), Propylene carbonate (PC), and N-Methyl pyrrolidone (NMP), further compositional analysis was performed using matrix-assisted laser desorption ionization-time of flight mass spectrometry (MALDI-TOF-MS) and solid-state nuclear magnetic resonance (ss-NMR). The MALDI-TOF-MS results ( Figure S13 ) indicated the presence of polymer components with >2000 Da in the SEI. Subsequently, ss-NMR dynamic-weighted analysis was performed to differentiate polymer and small molecule (non-polymer) components and estimate their relative contents. As shown in Figure 3f-h , the molar percentages of polymer in the F, B, and H spectra are estimated to be 79.5%, 69.7%, and 15.6%, respectively, indicating that most of the F and B are part of polymer components, with only 14.7% of F contributing to LiF. Calibration using 2,4,6-Trifluorophenylboronic acid (TFPBA) determined the molar ratio of F:B:H in the polymer to be approximately 13:6:1 ( Figure S14 ). Since B is exclusively from the salt and H is solely from the solvent, the formation of polymer components results from the interaction between the salt and the solvent in BAFF, wherein the contents of coordinated DFOB – anions and free-state NDFA solvents dominate in the solution structure (see Raman spectra in Figure S1c ). Additionally, the Li + conductivity in the SEI was estimated using electrochemical impedance spectroscopy (EIS, Figures 3i ), which showed the BAFF-derived SEI exhibits the lowest activation energy for Li + transport (37.4 kJ mol –1 ) compared to the reference electrolytes (39.8~53.3 kJ mol –1 ), indicating superior Li + conductivity for the former. Based on these results, the BAFF-derived SEI appears to have a single-layer amorphous configuration, with a B-F-based polymer as the dominant component, exhibiting sub-nanometer uniformity, excellent flexibility, and fast lithium-ion conductivity. These characteristics distinguish it from the “outer organic/inner inorganic” bilayer configuration derived from CME as well as the LiF-dominated inorganic mosaic configuration reported for AFE, LHCE, and HCDE. 36-38 Our previous study revealed that the SEI’s flexibility and the Cu substrate effect profoundly impact the performance of AFLMBs under high areal capacity of >5.0 mAh cm –2 16 . BAFF effectively mitigates the adverse substrate effect and renders a highly flexible polymer SEI capable of accommodating volume changes during cycling, thereby facilitating high-quality lithium deposition. However, extensive research indicates that, even if dense lithium deposition is achieved during the initial charge, maintaining its stability without a host during subsequent cycling remains challenging 16,39,40 . This difficulty arises from the irregular 3D dissolution of lithium grains during discharge: the 3D shrinkage of lithium grains produces numerous pores and increases the surface area of metallic lithium, exacerbating uneven dissolution via side reactions with the electrolyte; concurrently, the SEI undergoes contraction and aggregation, or even cracking, which causes a dynamically changing structure that deviates from the initial charged state 16 , thereby further deteriorating subsequent lithium deposition/dissolution (see schematical illustration in Figure 3j ). Theoretically, if the distribution of lithium-ion flux through SEI is sufficiently uniform, lithium dissolution would begin from the top surface of every lithium grain contacted with the electrolyte, following a 2D-dissolution manner. However, this rarely occurs in practice, likely due to the inhomogeneity of the SEI. As is well known, the SEI of lithium metal typically contains a variety of organic/inorganic components with different Li + conductivity and has a composite structure consisting of crystalline and amorphous nanoparticles of varying size 36,37,41,42 . Although an artificial SEI may be homogenous initially, it is difficult to maintain this homogeneity, as the interphase contacted with strongly reductive lithium metal evolves in both composition and structure during cycling 22,43,44 . Therefore, the SEI on lithium metal generally exhibits inhomogeneity at the microscale ( Figure 3b-e, S11 ), which could be the root cause of irregular 3D lithium dissolution. The BAFF-derived polymer-rich SEI, however, demonstrates a single amorphous layer with sub-nanometer uniformity and excellent Li + conductivity, thereby enabling 2D lithium dissolution. Moreover, due to its exceptional flexibility and lithiophilicity, the SEI remains tightly adhered to the lithium metal surface and evolves a self-adaptive mesh-film structure to accommodate the volume change throughout the dissolution process, maintaining the structural stability. Furthermore, after dissolution, the resulting mesh-structured SEI exhibits an ultra-thin face at the center of each mesh cell, which could serve as the preferred nucleation site for subsequent lithium deposition. Consequently, a highly reversible 2D lithium deposition/dissolution is achieved, overcoming the inherent shortcoming of lacking a stable host for lithium metal electrode (see Figure 3k ). Performances of practical 500 Wh kg –1 AFLMBs Figure 4a displays the cycling performance of home-made 2.7 Ah, 508 Wh kg –1 , 1668 Wh L –1 anode-free pouch cells under an external pressure of 200 kPa. Due to the high areal capacity of 5.6 mAh cm –2 , the current densities for 0.1C charging and 0.3C discharging correspond to 0.56 and 1.68 mA cm –2 , respectively. After 100 cycles of deep charge/discharge at 2.8~4.4 V (100% DoD), the capacity retention reached 80% with an average Coulombic efficiency of 99.6%, and no gas generation was observed ( Figure S15 ). Similar results were obtained for 6.4 Ah AFLMBs ( Figure S16 ). In contract, the lifespan of AFLMBs using reference electrolytes (AFE, LHCE, HCDE) are 18, 50, and 55 cycles, respectively. Among them, the HDCE-based AFLMB suffered from severe gas generation during cycling ( Figure S17 ). Increasing the discharge cutoff voltage from 2.8 V to 3.6 V allows for the release of 80% of the total capacity (80% DoD), resulting in a 2.16 Ah AFLMB with energy density of 406 Wh kg –1 and 1334 Wh L –1 . Under these conditions, the battery lifespan is remarkably improved, with cycle numbers of 250 and 300 corresponding to capacity retentions of 80% and 70%, respectively ( Figure 4b ). Moreover, the battery exhibits minimal polarization and no gas generation throughout the entire discharge process ( Figure S13 ). To the best of our knowledge, this represents the best performance reported for an AFLMB to date ( Figure S18, Table S7 ). To understand the capacity decay, the cycled batteries were disassembled. No significant change was observed in the morphology or reversible capacity of the cycled NCM811 cathode ( Figure S19 ), indicating that the AFLMB degradation primarily originates from the anode side. Mass spectrometry-D 2 O titration (MST) measurements were performed to quantitatively analyze the content of “dead lithium” and lithium hydride (LiH) on the Cu foils collected from fully discharged (to 1.0 V) AFLMBs with 20% capacity decay. As shown in Figure 4c , S20 , the contributions of dead Li and LiH to capacity loss in the retired BAFF battery (100% DoD) were only 3.5% and 0.43%, respectively, both of which are considerably smaller than those of the reference electrolytes (7.2~13.2%, 0.7~2.4%). The average growth rate of dead Li in BAFF was only 0.002 mAh cm –2 per cycle, which is less than 1/3 of LHCE (0.007 mAh cm –2 per cycle) and 1/20 of AFE (0.041 mAh cm –2 per cycle). For the cycling at 80% DoD, the growth rate of dead Li further halved to 0.001 mAh cm –2 per cycle. Hence, uneven lithium deposition is no longer the primary cause of capacity decay in the BAFF battery. Additionally, due to the minimal accumulation of chemically active dead Li and LiH, the BAFF battery should pose a lower safety risk compared to reference batteries. Generally, achieving high energy density in a battery compromises its power density. However, as shown in Figure 4d , our 2.7 Ah AFLMB delivered approximately 90% of its initial capacity at a current density of 5.6 mA cm –2 (1C), 41% at 28 mA cm –2 (5C), and even 24% at an ultra-high current density of 49.2 mA cm –2 (7C). With both high energy density and high-rate discharge capability, this BAFF battery achieves an ultra-high power density of 1998 W kg –1 (5C) at high energy density of 180 Wh kg –1 , 591 Wh L –1 or 2650 W kg –1 (7C) at 96 Wh kg –1 , 316Wh L –1 at the cell level, superior to state-of-the-art commercial supercapacitors and rechargeable batteries ( Figure 4e ). Additionally, due to the excellent ion-conductivity in BAFF and its-derived SEI over a wide temperature range, our AFLMB also shows excellent discharge performance at low temperatures, delivering approximately 83% and 66% of its room-temperature capacity at –20 and –40 ℃, respectively ( Figure 4d ). Finally, in terms of manufacturing cost, the absence of anode active materials together with a 100% capacity utilization in our AFLMBs could potentially reduce the cost per kWh by 15~25% compared to commercial graphite LIBs, assuming the electrolyte cost remains unchanged. Consequently, our AFLMBs offer ultra-high gravimetric/volumetric energy density, ultra-high power density output, wide temperature operation range, extended lifespan, as well as low cost, making them potentially suitable for various application scenarios. Conclusion In summary, we successfully designed and validated practical 500 Wh kg –1 -level AFLMBs in 2.7 Ah and 6.4 Ah pouch cells, demonstrating an enhanced lifespan through the use of the BAFF electrolyte. The BAFF-derived self-adaptive mesh-structured SEI ensures uniform ion flux and large-volume-change accommodation and supports lithium deposition and dissolution to occur in a highly reversible planar (2D) manner, in contrast to the conventional 3D manner, thereby overcoming the inherent structural instability of host-free lithium metal electrodes. As a result, the 2.7 Ah AFLMB (508 Wh kg –1 , 1668 Wh L –1 ) without any host-material coating achieved stable cycling performance for over 100 cycles at 100% DoD and 250 cycles at 80% DoD, maintaining 80% capacity retention, and delivered an unprecedented ultra-high power output of 2650 W kg –1 at 96 Wh kg –1 and 316 Wh L –1 . Compared to a best commercial graphite-based LIB (280 Wh kg –1 , 770 Wh L –1 ), our AFLMBs eliminate the need for active anode materials, resulting in an 81% increase in gravimetric energy density, a 117% increase in volumetric energy density, and a 15-25% reduction in cost per kWh. Given that a battery life of 300 cycles effectively meets the operational requirements of drones, AFLMBs hold substantial promise for applications in drone and electric aircraft markets, where high gravimetric and volumetric energy and power densities are critical. With further advancements in fast charging and lifespan, AFLMBs could potentially be extended to the electric vehicle market in the future. Methods Electrolyte preparation and their p hysic o chemical properties LiFSI, LiPF 6 , LiDFOB, LiBF 4 , DEC, FEC, FEMC, DME and commercial carbonate electrolyte ( 1.0 M LiPF 6 in EC/DMC ) were purchased from Duoduo Chem. HFE, TTE and NDFA was purchased from Tokyo Chemical Industry Co.. These solvents were dried by 3 Å molecular sieves for two days before use. The AFE electrolyte was prepared by dissolving 1.0 M LiPF 6 in FEC/FEMC/HFE (weight ratio=2:6:2). The LHCE electrolyte was prepared by dissolving 1.5 M LiFSI in DME/TTE (molar ratio=1.2:3). The HCDE electrolyte was prepared by dissolving 2.0 M LiDFOB and 1.4 M LiBF 4 in FEC/DEC (volume ratio=1:2). The BAFF electrolyte was prepared by dissolving 1.6 M LiDFOB in NDFA (molar ratio=1:5). Electrolytes were stored in an argon-filled glovebox (Mikrouna, oxygen <0.1 ppm, water <0.1 ppm) at room temperature. The water content of the solution, as measured by a Karl Fischer aquameter, was less 10 ppm. The ionic conductivity was measured by AC impedance spectrometer (Solartron, 1470E) in a symmetrical Pt|electrolyte|Pt cell. The viscosity and density of electrolytes were measured using a kinematic viscometer (Anton Paar, SVM 3001). The solution structure was studied by a Raman spectrometer (Anton Paar, Cora 5700) with an exciting laser of 785 nm. Anode-free pouch cells construction and e lectrochemical measurements Cu||NCM811 anode-free lithium pouch cell was assembled using double coated NCM811 (5.6 mAh cm – 2 per one side) as the cathode, commercial polyethylene (PE) as the separator, and bare Cu foil without any surface coating as the anode. All pouch cells were filled with 1.5 g Ah – 1 electrolyte and vacuum sealing in an argon glovebox. Detailed cell parameters are provided in Table S6 . All anode-free pouch cells were conducted under galvanostatic charge–discharge tests using NEWARE battery tester (NEWARE CT-4008-5V-6A). The constant-voltage charge process was applied until the charge current decayed to 0.05C. The cycling test was performed at 30 ℃ by charging at 0.1C and discharging at 0.3C, after two formation cycles at 0.1C. The temperature-dependent performance was tested by charging at 0.1C, 30 ℃ followed by discharging at different temperatures (-40 ~ 60 ℃). The rate performance was tested by charging at 0.1C, 30 ℃ followed by discharging at different rates using a NEWARE battery tester (CT-8008-5V60A-NTFA). A 1C rate corresponds to 5.6 mA cm – 2 . The resistances of Li + transport through SEIs were measured by a MPG-2 electrochemical workstation (Bio-Logic, France) in a symmetrical Li||Li cell. The Li electrodes were retracted from the charged Cu||NCM811 pouch cells after five cycles. Characterizations All the cycled samples were collected from pouch cells and rinsed by NDFA or DME solvent, followed by handling in transfer devices under argon atmosphere to avoid the air contamination. The morphologies of top, slope (30°) and bottom view of deposited Li samples were characterized using a field emission SEM (SU8230, Hitachi). The cross-sectional morphologies of deposited Li samples were characterized using an Cryo-FIB-SEM (Helios 5 UX, Thermo Fisher). The operating voltage and emission current of the electron beam were 5 kV and 0.2 nA, respectively. A gallium-ion beam source (30 kV) was used to mill the sample with 2 nA for pattern milling, 26 pA for imaging, and 0.2 nA for cross-section cleaning. The stage temperature was maintained at −195 °C to prevent beam damage.The crystallographic information of deposited Li samples after five cycles was characterized by 2D GIXD with Eiger2 R 500k 2D detector (D8 Discover, Bruker) with an incident angle of 0.2°. The microstructure of the BAFF-derived SEI on Li surface (after five cycles) was characterized by cryo-TEM (JEM 2100F, JEOL). The surface roughness and Young’s modulus of SEIs on Li surface were characterized by a Cypher ES instrument (OXFORD) using AM-FM module. The probe was calibrated by a standard sample of PVDF membrane with Young’s modulus of 2.45 GPa. The chemical composition of BAFF-derived SEIs on bare Cu foils (after five cycles) were characterized by XPS, MALDI-TOF-MS, and ss-NMR. XPS experiments were conducted on a X-ray photoelectron spectrometer with Al-Kα radiation (Escalab 250Xi, Thermo Fisher). A charge neutralizer was applied to compensate the sample surface charge. MALDI-TOF-MS measurements were performed on an AXIMA-Performance Shimadzu Mass Spectrometer in linear mode (Power: 110 W) using 2,5-dihydroxybenzoic acid (DHB) as matrix. Solid-state 1 H, 11 B and 19 F Magic Angle Spinning (MAS) NMR experiments were performed on a Bruker-AVANCE-500M NMR spectrometer using a 1.3 mm double-resonance MAS probe at 60 kHz spinning rate. Dynamic-weighted analysis was applied to differentiate the polymer and small molecule components using the Carr-Purcell-Meiboom-Gill (CPMG) scheme for T 2 -filtering. The amount of inactive metallic Li and LiH in the fully deposited samples were quantified by MST experiments on a Hiden Analytical Mass Spectrometer (QIC-20) according to previously reported procedures 16 . Declarations Data availability The data that support the findings of this study are available within this article and its Supplementary information. Additional data are available from the corresponding authors upon reasonable request. Acknowledgements This work was supported by Research Center for Industries of the Future (RCIF), Zhejiang Key Laboratory of 3D Micro/nano Fabrication and Characterization, Westlake Education Foundation, and National Natural Science Foundation of China (Grant No. 21975207). The authors thank Drs. Yangjian Lin and Qike Jiang from Instrumentation and Service Center for Physical Sciences at Westlake University for supporting in Cryo-FIB-SEM and Cryo-TEM characterizations and data interpretation, and Profs. Pengfei Hu, Zexin Jin and Jiuan Lv at Westlake University for discussion on mechanistic understanding. Besides, the authors thank Drs. Xiaohe Miao, Lin Liu, Pei Sheng and Ms. Huan Zhang (from Instrumentation and Service Center for Physical Sciences), Drs. Yinjuan Chen, Zhong Chen, Ms. Danyu Gu and Xin Li (from Instrumentation and Service Center for Molecular Science), Ms. Yingchun Wu (from Instrumentation and Service Center for 3D Micro/nano Fabrication) for their assistance in measurements at Instrumentation and Service Centers of Westlake University. Author contributions J.W. and L.L. designed the experiments. L.L. prepared the materials, performed the electrochemical measurements, and carried out SEM-EDS, AFM, XPS, and Raman characterizations. Y.X. guided L.L. to carry out the MST measurements. X.L. and Y.X. carried out the NMR analysis. J.W. and L.L. prepared the manuscript. J.W. conceived and directed the project. Competing interests J.W. and L.L. are inventors of the published patents (CN202210110653.X, PCT/CN2022/138887). References Armand, M. & Tarascon, J. M. Building better batteries. Nature 451 , 652-657 (2008). Lin, D. C., Liu, Y. & Cui, Y. Reviving the lithium metal anode for high-energy batteries. Nat. Nanotechnol. 12 , 194-206 (2017). Schmuch, R., Wagner, R., Horpel, G., Placke, T. & Winter, M. Performance and cost of materials for lithium-based rechargeable automotive batteries. Nat. Energy 3 , 267-278 (2018). Meng, Y. S., Srinivasan, V. & Xu, K. Designing better electrolytes. Science 378 , eabq3750 (2022). Liu, N. et al. A pomegranate-inspired nanoscale design for large-volume-change lithium battery anodes. Nat Nanotechnol. 9 , 187-192 (2014). Choi, S., Kwon, T. W., Coskun, A. & Choi, J. W. Highly elastic binders integrating polyrotaxanes for silicon microparticle anodes in lithium ion batteries. Science 357 , 279-283 (2017). Chen, Y. et al. Mechanical-electrochemical modeling of silicon-graphite composite anode for lithium-ion batteries. J. Power Sources 527 , 231178 (2022). Li, A. et al. Asymmetric electrolyte design for high-energy lithium-ion batteries with micro-sized alloying anodes. Nat. Energy , DOI:10.1038/s41560-024-01619-2 (2024). Cheng, X., Zhang, R., Zhao, C. & Zhang, Q. Toward safe lithium metal anode in rechargeable batteries: a review. Chem. Rev. 117 , 10403-10473 (2017). Liu, J. et al. Pathways for practical high-energy long-cycling lithium metal batteries. Nat. Energy 4 , 180-186 (2019). Chen, H. et al. Free-standing ultrathin lithium metal–graphene oxide host foils with controllable thickness for lithium batteries. Nat. Energy 6 , 790-798 (2021). Wu, W., Luo, W. & Huang, Y. Less is more: a perspective on thinning lithium metal towards high-energy-density rechargeable lithium batteries. Chem. Soc. Rev. 52 , 2553-2572 (2023). Nanda, S., Gupta, A. & Manthiram, A. Anode‐free full cells: a pathway to high‐energy density lithium‐metal batteries. Adv. Energy Mater. 11 , 2000804 (2020). Dong, L. et al. Toward practical anode-free lithium pouch batteries. Energy Environ. Sci. 16 , 5605-5632 (2023). Xu, P. et al. Anode‐free alkali metal batteries: from laboratory to practicability. Adv. Funct. Mater. , 2406080 (2024). Liu, L., Xiang, Y. & Wang, J. From cell to atomic level: understanding the degradation in 99% coulombic efficiency and 450 Wh kg⁻¹ anode-free pouch cells. ChemRxiv , DOI: 10.26434/chemrxiv-2024-pzdws (2024). Louli, A. J. et al. Diagnosing and correcting anode-free cell failure via electrolyte and morphological analysis. Nat. Energy 5 , 693-702 (2020). Qiao, Y. et al. A high-energy-density and long-life initial-anode-free lithium battery enabled by a Li 2 O sacrificial agent. Nat. Energy 6 , 653-662 (2021). Niu, C. et al. Balancing interfacial reactions to achieve long cycle life in high-energy lithium metal batteries. Nat. Energy 6 , 723-732 (2021). Yu, Z. et al. Rational solvent molecule tuning for high-performance lithium metal battery electrolytes. Nat. Energy 7 , 94-106 (2022). Wu, Z. et al. Deciphering and modulating energetics of solvation structure enables aggressive high-voltage chemistry of Li metal batteries. Chem 9 , 650-664 (2023). Wang, Y. et al. Anode-free lithium metal batteries based on an ultrathin and respirable interphase layer. Angew. Chem. Int. Ed. 62 , e202304978 (2023). Chen, L. et al. Enhancing the cycle-life of initial-anode-free lithium-metal batteries by pre-lithiation in Mn-based Li-rich spinel cathodes. J. Mater. Chem. A 11 , 11119-11125 (2023). Shi, J. et al. In situ p-block protective layer plating in carbonate-based electrolytes enables stable cell cycling in anode-free lithium batteries. Nat. Mater. , DOI:10.1038/s41563-024-01997-8 (2024). Zhang, X., Ma, L., Cai, Y., Fransaer, J. & Zheng, Q. A low-fermi-level current collector enables anode-free lithium metal batteries with long cycle life. Matter 7 , 583-602 (2024). Lin, L. et al. Li‐rich Li 2 [Ni 0.8 Co 0.1 Mn 0.1 ]O 2 for anode‐free lithium metal batteries. Angew. Chem. Int. Ed. 60 , 8289-8296 (2021). Lin, L. et al. Epitaxial induced plating current-collector lasting lifespan of anode-free lithium metal battery. Adv. Energy Mater. 11 , 2003709 (2021). Mao, M. et al. Anion-enrichment interface enables high-voltage anode-free lithium metal batteries. Nat. Commun. 14 , 1082 (2023). Liu, L. & Wang, J. Overcoming copper substrate thermodynamic limitations in anode-free lithium pouch cells via in situ seed implantation. Nano Lett. 23 , 10251-10258 (2023). Li, H. et al. Significance of current collectors for high performance conventional lithium‐ion batteries: a review. Adv. Funct. Mater. 33 , 2305515 (2023). Moon, C., Lian, J. & Lee, M. G. Identification of elastic and plastic properties of aluminum-polymer laminated pouch film for lithium-ion batteries: a hybrid experimental-numerical scheme. J. Energy Storage 72 , 108601 (2023). An, S. J. et al. Correlation of electrolyte volume and electrochemical performance in lithium-ion pouch cells with graphite anodes and NMC532 cathodes. J. Electrochem. Soc. 164 , A1195-A1202 (2017). An, S. J. et al. Electrolyte volume effects on electrochemical performance and solid electrolyte interphase in Si-graphite/NMC lithium-ion pouch cells. ACS Appl. Mater. Interfaces 9 , 18799-18808 (2017). Ue, M., Sakaushi, K. & Uosaki, K. Basic knowledge in battery research bridging the gap between academia and industry. Mater. Horiz. 7 , 1937-1954 (2020). Zhao, Q. et al. On the crystallography and reversibility of lithium electrodeposits at ultrahigh capacity. Nat. Commun. 12 , 6034 (2021). Xu, K. Interfaces and interphases in batteries. J. Power Sources 559 , 232652 (2023). Aurbach, D. Review of selected electrode–solution interactions which determine the performance of Li and Li ion batteries. J. Power Sources 89 , 206-218 (2000). Wang, J. et al. Fire-extinguishing organic electrolytes for safe batteries. Nat. Energy 3 , 22-29 (2018). Fang, C. et al. Pressure-tailored lithium deposition and dissolution in lithium metal batteries. Nat. Energy 6 , 987-994 (2021). Zhang, X. et al. Columnar lithium metal anodes. Angew. Chem. Int. Ed. 56 , 14207-14211 (2017). Weng, S. et al. Temperature-dependent interphase formation and Li + transport in lithium metal batteries. Nat. Commun. 14 , 4474 (2023). Li, Y. et al. Correlating structure and function of battery interphases at atomic resolution using cryoelectron microscopy. Joule 2 , 2167-2177 (2018). He, M., Guo, R., Hobold, G. M., Gao, H. & Gallant, B. M. The intrinsic behavior of lithium fluoride in solid electrolyte interphases on lithium. Proc. Natl. Acad. Sci. USA 117 , 73-79 (2020). Chen, H. et al. Uniform high ionic conducting lithium sulfide protection layer for stable lithium metal anode. Adv. Energy Mater. 9 , 1900858 (2019). Additional Declarations Yes there is potential Competing Interest. J.W. and L.L. are inventors of the published patents (CN202210110653.X, PCT/CN2022/138887). Supplementary Files SI20240906X.pdf Supporting information Cite Share Download PDF Status: Published Journal Publication published 17 Mar, 2026 Read the published version in Nature → Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. 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Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-5047161","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Physical Sciences - Article","associatedPublications":[],"authors":[{"id":381130834,"identity":"6c3fa149-2930-47a2-89da-015fb5db82d2","order_by":0,"name":"Jianhui Wang","email":"data:image/png;base64,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","orcid":"https://orcid.org/0000-0002-4170-1132","institution":"Westlake University","correspondingAuthor":true,"prefix":"","firstName":"Jianhui","middleName":"","lastName":"Wang","suffix":""},{"id":381130835,"identity":"5c616187-4861-4e28-8f87-11924fb39d6d","order_by":1,"name":"Lei Liu","email":"","orcid":"","institution":"Westlake University","correspondingAuthor":false,"prefix":"","firstName":"Lei","middleName":"","lastName":"Liu","suffix":""},{"id":381130836,"identity":"2753fec0-9597-4a6b-a0bc-09da0e3706ef","order_by":2,"name":"Yuxuan Xiang","email":"","orcid":"https://orcid.org/0000-0001-5741-1546","institution":"Westlake University","correspondingAuthor":false,"prefix":"","firstName":"Yuxuan","middleName":"","lastName":"Xiang","suffix":""},{"id":381130837,"identity":"3068f36b-87d7-47e5-ad06-3aa0bd10f916","order_by":3,"name":"Xingyu Lu","email":"","orcid":"","institution":"Westlake University","correspondingAuthor":false,"prefix":"","firstName":"Xingyu","middleName":"","lastName":"Lu","suffix":""}],"badges":[],"createdAt":"2024-09-07 05:21:17","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-5047161/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-5047161/v1","draftVersion":[],"editorialEvents":[{"content":"https://doi.org/10.1038/s41586-026-10402-0","type":"published","date":"2026-03-17T04:00:00+00:00"}],"editorialNote":"","failedWorkflow":false,"files":[{"id":73928823,"identity":"584501b2-21a9-4deb-a413-c32b21ef845a","added_by":"auto","created_at":"2025-01-16 05:36:48","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":939020,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eDesign of practical 500 Wh kg\u003c/strong\u003e\u003csup\u003e\u003cstrong\u003e–1\u003c/strong\u003e\u003c/sup\u003e\u003cstrong\u003e anode-free Li pouch cells.\u003c/strong\u003e (a) The trend of battery price dependent of energy density for LIBs, LMBs and AFLMBs. Detailed information is shown in Table S1. (b) Correlations of areal capacity, E/C ratio, and stack layer number for commercial 200, 268 Wh kg\u003csup\u003e–1\u003c/sup\u003e LIBs and 500 Wh kg\u003csup\u003e–1 \u003c/sup\u003eAFLMB. The LIBs adopt cell components (current collectors, separator, and Al-plastic film) same with LGES and Li-FUN products, respectively, while the AFLMB adopts commercially available, most lightweighted cell components. Detailed information and calculations are shown in Table S2-S4. (c) Schematic comparison in battery configuration of 2.7 Ah AFLMB and LIB using the same cell components. (d) Schematic illustration of the BAFF electrolyte. (e) Photograph of a home-made 2.7 Ah anode-free Cu||NCM811 pouch cell on the electronic balance.\u003c/p\u003e","description":"","filename":"1.png","url":"https://assets-eu.researchsquare.com/files/rs-5047161/v1/ac2261e7ce66670472bad4c3.png"},{"id":73928826,"identity":"d59f78c4-3c9c-40f8-ba22-ed9424b45bcd","added_by":"auto","created_at":"2025-01-16 05:36:48","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":1490166,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eMorphological and crystallographic evolution of 5.6 mAh cm\u003c/strong\u003e\u003csup\u003e\u003cstrong\u003e–1\u003c/strong\u003e\u003c/sup\u003e\u003cstrong\u003e Li deposition/dissolution. \u003c/strong\u003e(a) Initial five charge/discharge curves of anode-free Cu||NCM811 pouch cells using BAFF. (b) Cross-sectional Cryo-FIB-SEM images of 5th deposited Li at 4.4 V in BAFF. (c) 2D GIXD analysis of 5th deposited Li in BAFF (upper) and LHCE (lower). (d-h) Morphological changes of initial Li deposition/dissolution on bare Cu foil in BAFF: (d) 1st charge-3.6 V, (e) 1st charge-4.4 V, (f) 1st discharge-3.9 V, (g) 1st discharge-1.0 V and (h) 2nd charge-3.6 V. (i-j) EDS mapping images of Cu foils after full Li dissolution at 1.0 V: (i) BAFF: 1st and 5th; (j) LHCE: 1st and 2nd.\u003c/p\u003e","description":"","filename":"2.png","url":"https://assets-eu.researchsquare.com/files/rs-5047161/v1/3f21e2373ffe32dc82dd2505.png"},{"id":73928825,"identity":"7580a03d-2746-4db5-9257-f1fb895d5199","added_by":"auto","created_at":"2025-01-16 05:36:48","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":1148251,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eSEI Characterization and Li deposition/dissolution mechanism.\u003c/strong\u003e (a) Cryo-EM image of 5th deposited Li in BAFF. Inset is FFT analysis of SEI. (b-e) AFM results of 5th deposited Li in BAFF and LHCE: the height images (b) and their distribution (c); the Young’s modulus mapping (d) and their distribution (e). (f-h) ss-NMR spectra of BAFF-derived SEl collected at 5th deposition state: (f) \u003csup\u003e19\u003c/sup\u003eF, (g) \u003csup\u003e11\u003c/sup\u003eB, and (h) \u003csup\u003e1\u003c/sup\u003eH polymer and non-polymer species are differentiated by ss-NMR dynamic-weighted analysis. (i) Arrhenius plots for the resistance of Li\u003csup\u003e+\u003c/sup\u003e transport through different SEIs. (j,k) Schematic illustration of (j) 3D and (k) 2D Li deposition/dissolution dictated by SEI properties.\u003c/p\u003e","description":"","filename":"3.png","url":"https://assets-eu.researchsquare.com/files/rs-5047161/v1/029556e113764c5f008607f1.png"},{"id":73928824,"identity":"b7622a58-9fb2-46b1-9425-c39248a7ae2e","added_by":"auto","created_at":"2025-01-16 05:36:48","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":710030,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003ePerformances of home-made 2.7Ah, 508 Wh kg\u003c/strong\u003e\u003csup\u003e\u003cstrong\u003e–1\u003c/strong\u003e\u003c/sup\u003e\u003cstrong\u003e anode-free Cu||NCM811 pouch cells.\u003c/strong\u003e (a,b) Cycling performances of anode-free Li pouch cells using BAFF at (a) 100% DoD (2.8–4.4 V) and (b) 80% DoD (3.6–4.3 V). (c) MST-determined contributions of Dead Li and LiH to capacity loss in different electrolytes. (d) Discharge profiles of anode-free Li pouch cells using BAFF at different rates (upper) and different temperatures (lower). (e) The Ragone plot of home-made anode-free Li pouch cells and representative commercial energy storage devices. Detailed product information is provided in Table S8.\u003c/p\u003e","description":"","filename":"4.png","url":"https://assets-eu.researchsquare.com/files/rs-5047161/v1/a81a25dfd92990c98750b959.png"},{"id":107604006,"identity":"6a37a6ff-63ec-4897-81a6-872d1dccf7fa","added_by":"auto","created_at":"2026-04-23 07:16:15","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":5241994,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-5047161/v1/686ea7f7-7511-4864-b579-5650257eb3a2.pdf"},{"id":73928827,"identity":"580bdeed-514e-4165-ac48-f9354f52e8cb","added_by":"auto","created_at":"2025-01-16 05:36:49","extension":"pdf","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":12300494,"visible":true,"origin":"","legend":"Supporting information","description":"","filename":"SI20240906X.pdf","url":"https://assets-eu.researchsquare.com/files/rs-5047161/v1/c133322f1ef902b4f46fc792.pdf"}],"financialInterests":"\u003cb\u003eYes\u003c/b\u003e there is potential Competing Interest.\nJ.W. and L.L. are inventors of the published patents (CN202210110653.X, PCT/CN2022/138887).","formattedTitle":"Planar lithium deposition/dissolution enabling practical 500 Wh kg–1 anode-free pouch cells","fulltext":[{"header":"Introduction","content":"\u003cp\u003eThe quest for high-energy-density and low-cost rechargeable batteries has been a long-standing goal. After more than 30 years of commercial development, the energy density and cost of traditional lithium-ion batteries (LIBs) have nearly reached their theoretical limits\u003csup\u003e1-4\u003c/sup\u003e. The use of silicon (Si) to partially replace the graphite anode material in silicon-carbon (Si/C) composites can exceed the theoretical energy density limit of 280 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e for traditional LIBs. However, Si inherently suffers from poor electronic conductivity, low initial Coulombic efficiency, large volumetric expansion, and severe pulverization during cycling. These issues necessitate complex material engineering processes for practical application of high-specific-capacity Si/C anodes\u003csup\u003e5-8\u003c/sup\u003e. Consequently, the production cost of 360 Wh kg\u003csup\u003e\u0026ndash;1\u0026nbsp;\u003c/sup\u003eSi-based LIBs is nearly double that of traditional LIBs, and mass production has yet to be realized.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eThe lithium metal batteries (LMBs), which utilize lithium metal as the anode and high-nickel nickel-cobalt-manganese (NCM) as the cathode, offer an energy density exceeding 400 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e\u003csup\u003e9,10\u003c/sup\u003e. However, the high chemical reactivity of metallic lithium and the challenges in processing ultra-thin lithium foil pose significant obstacles to the large-scale application of LMBs\u003csup\u003e11,12\u003c/sup\u003e, particularly in terms of production costs, safety, and product consistency. Moreover, both Si-based LIBs and LMBs are likely to experience a more pronounced increase in production costs as energy density rises (see \u003cstrong\u003eFigure 1a\u003c/strong\u003e).\u003c/p\u003e\n\u003cp\u003eIn contrast, anode-free lithium metal batteries (AFLMBs), which do not use anode active materials during manufacturing, simplify battery configuration and reduce raw material usage\u003csup\u003e13-15\u003c/sup\u003e. This battery design offers the highest possible gravimetric and volumetric energy densities while reducing the production cost per unit of energy to a level even lower than that of traditional LIBs. Moreover, the cost of production does not significantly increase with higher energy density, positioning AFLMBs as an ideal approach for mass-producing high-energy-density batteries (\u003cstrong\u003eFigure 1a\u003c/strong\u003e). However, the lack of an anode host and excess lithium resource subjects AFLMBs to extremely harsh operating conditions, resulting in a shortened lifespan, especially under high-energy-density conditions\u003csup\u003e14,16\u003c/sup\u003e. To date, most AFLMBs have demonstrated limited cycling life (50-300 cycles) at relatively low energy densities (50-350 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e) \u003csup\u003e17-25\u003c/sup\u003e, rendering them less competitive compared to traditional LIBs (160-280 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, \u0026gt;800 cycles). Several attempts to develop AFLMBs at the 400 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e level have all showed a short lifespan of less than 80 cycles\u003csup\u003e16,26-29\u003c/sup\u003e. Currently, there have been no reports of repeatedly rechargeable AFLMBs achieving energy densities exceeding 500 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eDesign of practical 500 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e pouch cells\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eMaximizing the active electrode materials while minimizing inactive materials is essential for constructing high-energy-density lithium batteries. In an AFLMB with a given cathode material, the energy density is determined by the cathode areal capacity, the number of stack layers, and the electrolyte weight to cell capacity (E/C), assuming that other technological parameters, such as current collectors, separator, tabs, and Al-plastic film, are established. The use of high areal capacity, multiple layers, and low E/C enhances battery energy density. However, excessive areal capacity and stack layers, or too low E/C, can introduce challenges in battery manufacturing and performances. Currently, commercial Ah-level LIBs with graphite anodes and NCM cathodes typically feature cell parameters of 5 ~ 20 cathode layers, 2.0 ~ 5.0 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e areal capacity, and 1.5 ~ 4.0 g Ah\u003csup\u003e\u0026ndash;1\u003c/sup\u003e E/C for optimal battery performances. Moreover, the current collectors for the anode (Cu foil, 4~12 \u0026mu;m thickness) and cathode (Al foil, 12~20 \u0026mu;m thickness), as well as the Al-plastic film (\u0026gt;88 \u0026mu;m thickness), are indispensable due to their chemical/mechanical stability and cost-effectiveness in Ah-level practical applications\u003csup\u003e30,31\u003c/sup\u003e. Therefore, to achieve the highest possible energy density under practical conditions, we selected 4.4 V NCM811 as the cathode material, owing to its highest specific energy density (216 mAh g\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, 3.84 V average discharge voltage) among currently mature cathode materials, and minimized the weight of inactive materials by using 4 \u0026mu;m Cu foil, 12 \u0026mu;m Al foil, 12 \u0026mu;m polyethylene (PE) separator, 100 \u0026mu;m tabs, and 88 \u0026mu;m Al-plastic film.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eA mathematical model of pouch cells was established to correlate the cell energy density with key parameters of the cathode areal capacity, stack layer number, and E/C. The detailed calculations are provided in\u003cstrong\u003e\u0026nbsp;Table S2\u0026ndash;S4\u003c/strong\u003e. As shown in \u003cstrong\u003eFigure 1b\u003c/strong\u003e, our calculated correlation matches well with the disclosed cell parameters of commercial LIBs (LGES, 268 Wh kg\u003csup\u003e\u0026ndash;1\u0026nbsp;\u003c/sup\u003egraphite||NCM712 and Li-FUN, 200 Wh kg\u003csup\u003e\u0026ndash;1\u0026nbsp;\u003c/sup\u003egraphite||NCM523), validating the accuracy of our model. According to the calculated AFLMB correlation, constructing a 500 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e AFLMB has a considerably higher requirement as compared to commercial LIBs, despite the absence of anode material. If E/C \u0026gt; 2.0 g Ah\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, it requires more than 20 cathode layers with an areal capacity exceeding 6.8 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e, which is too high as compared to commercial NCM cathodes (2 ~ 5 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e). If the cathode areal capacity is less than 4.0 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e, the E/C ratio would need to be below 1.3 g Ah\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, which hardly ensures adequate wetting and rate capability regarding the porosity of electrodes and separator\u003csup\u003e32-34\u003c/sup\u003e. Furthermore, if the stack comprises fewer than 10 cathode layers, the weight percentage of Al-plastic film in the cell would be too high to reach the target. Considering all these factors, we ultimately selected cell parameters of 5.6 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e areal capacity, 10 stack layers, and 1.5 g Ah\u003csup\u003e\u0026ndash;1\u003c/sup\u003e E/C to construct a 500 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e-level AFLMB (\u003cstrong\u003eFigure 1\u003c/strong\u003e\u003cstrong\u003eb, c\u003c/strong\u003e). These parameters are notably more stringent than those in previously reported AFLMBs (2 ~ 4 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e, 2 stack layers, and \u0026gt; 2.0 g Ah\u003csup\u003e\u0026ndash;1\u003c/sup\u003e)\u0026nbsp;\u003csup\u003e16-29\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003eThe E/C ratio of 1.5 g Ah\u003csup\u003e\u0026ndash;1\u003c/sup\u003e approaches the lower limit for advanced LIBs\u003csup\u003e32-34\u003c/sup\u003e, imposing rigorous demands on the density, wettability, and stability of the electrolyte. Currently, only liquid electrolytes could meet these requirements. In this study, we selected an optimized electrolyte of 1.6 M lithium difluoro(oxalate)borate in N,N-Dimethyltrifluoroacetamide solvent (1.6 M LiDFOB/NDFA, named as BAFF electrolyte, illustrated in \u003cstrong\u003eFigure 1d\u003c/strong\u003e) for the following reasons: (1) the electrolyte has a low density (1.36 g cm\u003csup\u003e\u0026ndash;3\u003c/sup\u003e), viscosity (7.04 mPa s) and vapor pressure (1.15 kPa), as well as good wettability with cell materials at room temperature (\u003cstrong\u003eFigure S1a, Table S5\u003c/strong\u003e); (2) it exhibits high ion conductivities across a wide temperature range of \u0026ndash;40 to +60 ℃ (0.5 ~ 14.2 mS cm\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, \u003cstrong\u003eFigure S1b\u003c/strong\u003e); (3) it has a wide electrochemical stability window of ~5.0 V, enabling not only stable cycling of NCM811 to 4.5 V (\u003cstrong\u003eFigure S2-S3\u003c/strong\u003e), but also highly reversible planar Li deposition/dissolution of 5.6 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e, thereby enhancing the longevity of 500 Wh kg\u003csup\u003e\u0026ndash;1\u0026nbsp;\u003c/sup\u003eAFLMB (see discussed below); and (4) it produces almost no gas during cycling, overcoming the gassing issue generally faced by LiDFOB-based electrolytes (\u003cstrong\u003eFigure S3\u003c/strong\u003e)\u003csup\u003e17,28\u003c/sup\u003e. Additionally, we tested AFLMBs using three advanced electrolytes reported in literature, including local high concentration electrolyte (LHCE), high concentration dual-salt electrolyte (HCDE), and all fluorinated electrolyte (AFE), for comparison. As depicted in \u003cstrong\u003eFigure 1c-e\u003c/strong\u003e, our designed AFLMB exhibits a total capacity of 2.7 Ah and total mass of 20.5 g, achieving a gravimetric energy density of 508 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e (\u003cstrong\u003eTable S6\u003c/strong\u003e). Moreover, the absence of low-density graphite materials significantly reduces the cell thickness from 4.3 mm to 2.3 mm, resulting in an ultra-high volumetric energy density of 1668 Wh L\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, which is more than double that of the LIB using the same cell components (\u003cstrong\u003eFigure 1c\u003c/strong\u003e). When excluding the Al-plastic film, the stack energy densities of our AFLMB reach 543 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e and 1806 Wh L\u003csup\u003e\u0026ndash;1\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eReversible planar lithium-orientated deposition/dissolution\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eAlthough AFLMBs do not utilize active anode materials during manufacturing, a micrometer-scale layer of lithium metal is deposited on the anode current collector, whose stability during cycling dictates the battery lifespan. \u003cstrong\u003eFigure 2a\u003c/strong\u003e illustrates the initial five cycles of a Cu||NCM811 anode-free pouch cell using the BAFF electrolyte at a current density of 0.56 mA cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e. The Coulombic efficiency is 93.8% in the first cycle and rapidly increases to 99.8% in the third cycle, indicating highly reversible lithium deposition/dissolution under the high areal capacity of 5.6 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e. A reversible specific capacity of 216 mAh g\u003csup\u003e\u0026ndash;1\u003c/sup\u003e was obtained in the pouch cell, which is equal to that of Li||NCM811 half-cell, indicating a 100% capacity utilization. In sharp contrast, dendritic short-circuiting occurs in the conventional carbonate electrolyte (CME) even during the first charge under these harsh conditions (see \u003cstrong\u003eFigure S4a\u003c/strong\u003e). To evaluate the quality of the deposited lithium on the bare Cu current collector, charged samples were collected after five cycles and analyzed by cryogenic focused ion beam scanning electron microscope (Cryo-FIB-SEM). The cross-sectional image (\u003cstrong\u003eFigure 2b\u003c/strong\u003e) shows a deposited lithium thickness of 28 \u0026mu;m, which closely matches the theoretic thickness of 27 \u0026mu;m for 5.6 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e of metallic lithium. The upper half of the deposited layer is composed of columnar grains approximately 8 \u0026mu;m in diameter, while the lower part exhibits no grain boundaries, demonstrating an integrated and compact structure. Through a substantial number of cross-sectional slice images, a 3D reconstruction of the deposited lithium with dimensions 10\u0026times;28\u0026times;5 \u0026mu;m reveals a porosity of only 1.6% for BAFF, which is remarkably superior to those obtained with reference electrolytes AFE, LHCE and HCDE, whose deposited lithium thicknesses are 37 \u0026mu;m, 35 \u0026mu;m, and 33 \u0026mu;m with corresponding porosities of 20.0%, 16.6%, and 11.1%, respectively (see \u003cstrong\u003eFigure S5\u003c/strong\u003e).\u003c/p\u003e\n\u003cp\u003eSubsequently, the deposited lithium was subjected to\u0026nbsp;2D grazing-incidence X-ray diffraction measurement (2D GIXD). As shown in \u003cstrong\u003eFigure 2c\u003c/strong\u003e-upper, distinct flecks of the (110) crystal plane can be clearly observed for BAFF, demonstrating a Li\u003csub\u003e(110)\u003c/sub\u003e texture characteristic, which indicates that BAFF enables the orientated growth of metallic lithium on bare Cu substrate. This finding contrasts sharply with reference electrolytes (see \u003cstrong\u003eFigure 2c\u003c/strong\u003e-lower, \u003cstrong\u003eS6\u003c/strong\u003e) and prior studies\u003csup\u003e16,29\u003c/sup\u003e, wherein abundant lithium mosses typically grow on the Cu surface (\u003cstrong\u003eFigure S7a\u003c/strong\u003e) due to its lithiophobic and uneven nature. To induce lithium orientated growth on bare Cu substrate, it is necessary to firstly deposit lithium in a thickness \u0026gt;50 \u0026mu;m to overcome the adverse substrate effect\u003csup\u003e35\u003c/sup\u003e. Surprisingly, the use of BAFF directly grows lithium texture on bare Cu with minimal nucleation overpotential (8 mV, \u003cstrong\u003eFigure 2a inset, S4c\u003c/strong\u003e) and lithium mosses (see bottom view in \u003cstrong\u003eFigure S7a, S9a\u003c/strong\u003e), indicating that BAFF effectively mitigates the adverse impact of the Cu substrate.\u003c/p\u003e\n\u003cp\u003eFurther, scanning electron microscope (SEM) was employed to directly observe the morphology changes of deposited lithium at different states during cycling. In BAFF, initial charge to 3.6 V produced a uniform distribution of lithium nuclei with diameter of approximately 200 nm on bare Cu (\u003cstrong\u003eFigure 2d\u003c/strong\u003e), which grew into large grains with sizes of 5~8 \u0026mu;m upon charging to 4.4 V (\u003cstrong\u003eFigure 2e\u003c/strong\u003e). Interestingly, in the following discharge to 3.9 V, the dissolution of lithium resulted in a \u0026ldquo;hemispherical pit\u0026rdquo; on each grain (\u003cstrong\u003eFigure 2f\u003c/strong\u003e). Since the lithium metal surface is covered by a solid electrolyte interphase (SEI), the deep indentation morphology reflects that the SEI keeps firmly adhered to the shrinking lithium metal without fracturing, demonstrating exceptional flexibility and lithiophilicity of the SEI. After fully discharging to 1.0 V, all the deposited lithium metal dissolved, leaving behind a unique mesh film covering the Cu surface (\u003cstrong\u003eFigure 2g\u003c/strong\u003e). This is attributed to SEI aggregation caused by the full stripping of lithium columnar grains: the vertical contraction of SEI on the side surface of columnar grains forms the thick wall of mesh cell, while the horizontal downshift of SEI on the upper surface of columnar grains forms ultra-thin face of mesh cell, beneath which the Cu substrate can be observed via energy-dispersive spectroscopy (EDS) mapping (\u003cstrong\u003eFigure 2i, S7, S9\u003c/strong\u003e). Additionally, the size of these mesh cells (5~8 \u0026mu;m) is consistent with that of lithium columnar grains, suggesting minimal change in lithium grain size during the dissolution process. Upon the second charge to 3.6 V, this SEI-aggregated mesh film re-bulged as lithium metal plating (\u003cstrong\u003eFigure 2h\u003c/strong\u003e), indicating a reversible and self-adaptive process. Notably, throughout the charge/discharge processes, all lithium grains exhibit nearly synchronous bulge/pit surface, demonstrating a typical planar (2D) deposition/dissolution characteristic (\u003cstrong\u003eFigure S8, S10\u003c/strong\u003e). Meanwhile, the SEI remains firmly adhered to the lithium or Cu surface, effectively preventing electrolyte infiltration (\u003cstrong\u003eFigure 2g\u003c/strong\u003e). Both collectively ensure maximum structural stability during cycling, as strongly supported by EDS results (\u003cstrong\u003eFigure 2i\u003c/strong\u003e), which show a clear mesh-structure film on bare Cu for many cycles. By contrast, lithium deposition/dissolution in all the reference electrolytes shows very different behaviors (\u003cstrong\u003eFigure 2j, S7, S9\u003c/strong\u003e). During discharge, lithium grains contract and the gaps between grains increase following a typical 3D dissolution process, leading to either porous SEI (e.g., LHCE and HCDE) or cracked SEI (e.g., AFE) (\u003cstrong\u003eFigure S7\u003c/strong\u003e). This loose structure cannot prevent electrolyte infiltration, which inevitably creates numerous new sites for lithium nucleation and growth in subsequent charging, resulting in dynamically changing structure during cycling (\u003cstrong\u003eFigure S9\u003c/strong\u003e). Consequently, the planar lithium deposition/dissolution unique to BAFF renders a stable mesh-film-structure SEI on lithium metal during cycling, effectively overcoming its inherent shortcoming of lacking a host.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eSEI characterization and mechanistic understanding\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eTo understand the mechanism behind the unique planar lithium deposition/dissolution, we conducted a detailed study of the composition, structure, and mechanical and ion-conductive properties of the SEI after five cycles. As shown in \u003cstrong\u003eFigure 3a\u003c/strong\u003e, cryogenic transmission electron microscopy (Cryo-TEM) revealed that the BAFF-derived SEI on the Li\u003csub\u003e(110)\u003c/sub\u003e grains is approximately 8 nm thick and in amorphous state (see fast Fourier transform (FFT) analysis in \u003cstrong\u003eFigure 3a insert\u003c/strong\u003e). The nano-scale morphology and mechanical properties of the SEI were characterized by atomic force microscope (AFM). \u003cstrong\u003eFigures 3b-e\u003c/strong\u003e and \u003cstrong\u003eS11\u003c/strong\u003e show that the SEI has an exceptionally smooth surface with an ultra-low roughness (Ra = 0.7 nm) and a relatively high Young\u0026rsquo;s modulus of 3.15 GPa with a very narrow distribution (Ra = 0.21 GPa). This character of ultra-high uniformity at sub-nanometer scale for the BAFF-derived SEI is in sharp contrast to the significant heterogeneity observed in the SEIs from the reference electrolytes (LHCE, HCDE, and AFE), which exhibited considerably rougher surface (Ra = 4.6, 3.3, and 9.2 nm, respectively) and more widely distributed Young\u0026rsquo;s modulus values (Ra = 0.75, 0.52 and 0.82 GPa, respectively).\u003c/p\u003e\n\u003cp\u003eRegarding to the chemical composition of the BAFF-derived SEI, X-ray photoelectron spectroscopy (XPS) identified the primary elements (B, C, N, O and F) and some ingredients like organic B-F or C-F species, B-O species, and a small amount of LiF (\u003cstrong\u003eFigure S12\u003c/strong\u003e). Due to the SEI\u0026rsquo;s susceptibility to air and its insolubility in various high-polarity aprotic solvents such as Dimethyl sulfoxide (DMSO), Propylene carbonate (PC), and N-Methyl pyrrolidone (NMP), further compositional analysis was performed using matrix-assisted laser desorption ionization-time of flight mass spectrometry (MALDI-TOF-MS) and solid-state nuclear magnetic resonance (ss-NMR). The MALDI-TOF-MS results (\u003cstrong\u003eFigure S13\u003c/strong\u003e) indicated the presence of polymer components with \u0026gt;2000 Da in the SEI. Subsequently, ss-NMR dynamic-weighted analysis was performed to differentiate polymer and small molecule (non-polymer) components and estimate their relative contents. As shown in \u003cstrong\u003eFigure 3f-h\u003c/strong\u003e, the molar percentages of polymer in the F, B, and H spectra are estimated to be 79.5%, 69.7%, and 15.6%, respectively, indicating that most of the F and B are part of polymer components, with only 14.7% of F contributing to LiF. Calibration using 2,4,6-Trifluorophenylboronic acid (TFPBA) determined the molar ratio of F:B:H in the polymer to be approximately 13:6:1 (\u003cstrong\u003eFigure S14\u003c/strong\u003e). Since B is exclusively from the salt and H is solely from the solvent, the formation of polymer components results from the interaction between the salt and the solvent in BAFF, wherein the contents of coordinated DFOB\u003csup\u003e\u0026ndash;\u003c/sup\u003e anions and free-state NDFA solvents dominate in the solution structure (see Raman spectra in \u003cstrong\u003eFigure S1c\u003c/strong\u003e). Additionally, the Li\u003csup\u003e+\u003c/sup\u003e conductivity in the SEI was estimated using electrochemical impedance spectroscopy (EIS, \u003cstrong\u003eFigures 3i\u003c/strong\u003e), which showed the BAFF-derived SEI exhibits the lowest activation energy for Li\u003csup\u003e+\u003c/sup\u003e transport (37.4 kJ mol\u003csup\u003e\u0026ndash;1\u003c/sup\u003e) compared to the reference electrolytes (39.8~53.3 kJ mol\u003csup\u003e\u0026ndash;1\u003c/sup\u003e), indicating superior Li\u003csup\u003e+\u003c/sup\u003e conductivity for the former. Based on these results, the BAFF-derived SEI appears to have a single-layer amorphous configuration, with a B-F-based polymer as the dominant component, exhibiting sub-nanometer uniformity, excellent flexibility, and fast lithium-ion conductivity. These characteristics distinguish it from the \u0026ldquo;outer organic/inner inorganic\u0026rdquo; bilayer configuration derived from CME as well as the LiF-dominated inorganic mosaic configuration reported for AFE, LHCE, and HCDE.\u003csup\u003e36-38\u003c/sup\u003e\u003c/p\u003e\n\u003cp\u003eOur previous study revealed that the SEI\u0026rsquo;s flexibility and the Cu substrate effect profoundly impact the performance of AFLMBs under high areal capacity of \u0026gt;5.0 mAh cm\u003csup\u003e\u0026ndash;2\u0026nbsp;\u003c/sup\u003e\u003csup\u003e16\u003c/sup\u003e. BAFF effectively mitigates the adverse substrate effect and renders a highly flexible polymer SEI capable of accommodating volume changes during cycling, thereby facilitating high-quality lithium deposition. However, extensive research indicates that, even if dense lithium deposition is achieved during the initial charge, maintaining its stability without a host during subsequent cycling remains challenging\u003csup\u003e16,39,40\u003c/sup\u003e. This difficulty arises from the irregular 3D dissolution of lithium grains during discharge: the 3D shrinkage of lithium grains produces numerous pores and increases the surface area of metallic lithium, exacerbating uneven dissolution via side reactions with the electrolyte; concurrently, the SEI undergoes contraction and aggregation, or even cracking, which causes a dynamically changing structure that deviates from the initial charged state\u003csup\u003e16\u003c/sup\u003e, thereby further deteriorating subsequent lithium deposition/dissolution (see schematical illustration in \u003cstrong\u003eFigure 3j\u003c/strong\u003e).\u003c/p\u003e\n\u003cp\u003eTheoretically, if the distribution of lithium-ion flux through SEI is sufficiently uniform, lithium dissolution would begin from the top surface of every lithium grain contacted with the electrolyte, following a 2D-dissolution manner. However, this rarely occurs in practice, likely due to the inhomogeneity of the SEI. As is well known, the SEI of lithium metal typically contains a variety of organic/inorganic components with different Li\u003csup\u003e+\u003c/sup\u003e conductivity and has a composite structure consisting of crystalline and amorphous nanoparticles of varying size\u003csup\u003e36,37,41,42\u003c/sup\u003e. Although an artificial SEI may be homogenous initially, it is difficult to maintain this homogeneity, as the interphase contacted with strongly reductive lithium metal evolves in both composition and structure during cycling\u003csup\u003e22,43,44\u003c/sup\u003e. Therefore, the SEI on lithium metal generally exhibits inhomogeneity at the microscale (\u003cstrong\u003eFigure 3b-e, S11\u003c/strong\u003e), which could be the root cause of irregular 3D lithium dissolution. The BAFF-derived polymer-rich SEI, however, demonstrates a single amorphous layer with sub-nanometer uniformity and excellent Li\u003csup\u003e+\u003c/sup\u003e conductivity, thereby enabling 2D lithium dissolution. Moreover, due to its exceptional flexibility and lithiophilicity, the SEI remains tightly adhered to the lithium metal surface and evolves a self-adaptive mesh-film structure to accommodate the volume change throughout the dissolution process, maintaining the structural stability. Furthermore, after dissolution, the resulting mesh-structured SEI exhibits an ultra-thin face at the center of each mesh cell, which could serve as the preferred nucleation site for subsequent lithium deposition. Consequently, a highly reversible 2D lithium deposition/dissolution is achieved, overcoming the inherent shortcoming of lacking a stable host for lithium metal electrode (see \u003cstrong\u003eFigure 3k\u003c/strong\u003e).\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003ePerformances of practical 500 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e AFLMBs\u0026nbsp;\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eFigure 4a\u003c/strong\u003e displays the cycling performance of home-made 2.7 Ah, 508 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, 1668 Wh L\u003csup\u003e\u0026ndash;1\u003c/sup\u003e anode-free pouch cells under an external pressure of 200 kPa. Due to the high areal capacity of 5.6 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e, the current densities for 0.1C charging and 0.3C discharging correspond to 0.56 and 1.68 mA cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e, respectively. After 100 cycles of deep charge/discharge at 2.8~4.4 V (100% DoD), the capacity retention reached 80% with an average Coulombic efficiency of 99.6%, and no gas generation was observed (\u003cstrong\u003eFigure S15\u003c/strong\u003e). Similar results were obtained for 6.4 Ah AFLMBs (\u003cstrong\u003eFigure S16\u003c/strong\u003e). In contract, the lifespan of AFLMBs using reference electrolytes (AFE, LHCE, HCDE) are 18, 50, and 55 cycles, respectively. Among them, the HDCE-based AFLMB suffered from severe gas generation during cycling (\u003cstrong\u003eFigure S17\u003c/strong\u003e). Increasing the discharge cutoff voltage from 2.8 V to 3.6 V allows for the release of 80% of the total capacity (80% DoD), resulting in a 2.16 Ah AFLMB with energy density of 406 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e and 1334 Wh L\u003csup\u003e\u0026ndash;1\u003c/sup\u003e. Under these conditions, the battery lifespan is remarkably improved, with cycle numbers of 250 and 300 corresponding to capacity retentions of 80% and 70%, respectively (\u003cstrong\u003eFigure 4b\u003c/strong\u003e). Moreover, the battery exhibits minimal polarization and no gas generation throughout the entire discharge process (\u003cstrong\u003eFigure S13\u003c/strong\u003e). To the best of our knowledge, this represents the best performance reported for an AFLMB to date (\u003cstrong\u003eFigure S18, Table S7\u003c/strong\u003e).\u003c/p\u003e\n\u003cp\u003eTo understand the capacity decay, the cycled batteries were disassembled. No significant change was observed in the morphology or reversible capacity of the cycled NCM811 cathode (\u003cstrong\u003eFigure S19\u003c/strong\u003e), indicating that the AFLMB degradation primarily originates from the anode side. Mass spectrometry-D\u003csub\u003e2\u003c/sub\u003eO titration (MST) measurements were performed to quantitatively analyze the content of \u0026ldquo;dead lithium\u0026rdquo; and lithium hydride (LiH) on the Cu foils collected from fully discharged (to 1.0 V) AFLMBs with 20% capacity decay. As shown in \u003cstrong\u003eFigure 4c\u003c/strong\u003e, \u003cstrong\u003eS20\u003c/strong\u003e, the contributions of dead Li and LiH to capacity loss in the retired BAFF battery (100% DoD) were only 3.5% and 0.43%, respectively, both of which are considerably smaller than those of the reference electrolytes (7.2~13.2%, 0.7~2.4%). The average growth rate of dead Li in BAFF was only 0.002 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e per cycle, which is less than 1/3 of LHCE (0.007 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e per cycle) and 1/20 of AFE (0.041 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e per cycle). For the cycling at 80% DoD, the growth rate of dead Li further halved to 0.001 mAh cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e per cycle. Hence, uneven lithium deposition is no longer the primary cause of capacity decay in the BAFF battery. Additionally, due to the minimal accumulation of chemically active dead Li and LiH, the BAFF battery should pose a lower safety risk compared to reference batteries.\u003c/p\u003e\n\u003cp\u003eGenerally, achieving high energy density in a battery compromises its power density. However, as shown in \u003cstrong\u003eFigure 4d\u003c/strong\u003e, our 2.7 Ah AFLMB delivered approximately 90% of its initial capacity at a current density of 5.6 mA cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e (1C), 41% at 28 mA cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e (5C), and even 24% at an ultra-high current density of 49.2 mA cm\u003csup\u003e\u0026ndash;2\u003c/sup\u003e (7C). With both high energy density and high-rate discharge capability, this BAFF battery achieves an ultra-high power density of 1998 W kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e (5C) at high energy density of 180 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, 591 Wh L\u003csup\u003e\u0026ndash;1\u003c/sup\u003e or 2650 W kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e (7C) at 96 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, 316Wh L\u003csup\u003e\u0026ndash;1\u003c/sup\u003e at the cell level, superior to state-of-the-art commercial supercapacitors and rechargeable batteries (\u003cstrong\u003eFigure 4e\u003c/strong\u003e). Additionally, due to the excellent ion-conductivity in BAFF and its-derived SEI over a wide temperature range, our AFLMB also shows excellent discharge performance at low temperatures, delivering approximately 83% and 66% of its room-temperature capacity at \u0026ndash;20 and \u0026ndash;40 ℃, respectively (\u003cstrong\u003eFigure 4d\u003c/strong\u003e).\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eFinally, in terms of manufacturing cost, the absence of anode active materials together with a 100% capacity utilization in our AFLMBs could potentially reduce the cost per kWh by 15~25% compared to commercial graphite LIBs, assuming the electrolyte cost remains unchanged. Consequently, our AFLMBs offer ultra-high gravimetric/volumetric energy density, ultra-high power density output, wide temperature operation range, extended lifespan, as well as low cost, making them potentially suitable for various application scenarios.\u003c/p\u003e"},{"header":"Conclusion","content":"\u003cp\u003eIn summary, we successfully designed and validated practical 500 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e-level AFLMBs in 2.7 Ah and 6.4 Ah pouch cells, demonstrating an enhanced lifespan through the use of the BAFF electrolyte. The BAFF-derived self-adaptive mesh-structured SEI ensures uniform ion flux and large-volume-change accommodation and supports lithium deposition and dissolution to occur in a highly reversible planar (2D) manner, in contrast to the conventional 3D manner, thereby overcoming the inherent structural instability of host-free lithium metal electrodes. As a result, the 2.7 Ah AFLMB (508 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, 1668 Wh L\u003csup\u003e\u0026ndash;1\u003c/sup\u003e) without any host-material coating achieved stable cycling performance for over 100 cycles at 100% DoD and 250 cycles at 80% DoD, maintaining 80% capacity retention, and delivered an unprecedented ultra-high power output of 2650 W kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e at 96 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e and 316 Wh L\u003csup\u003e\u0026ndash;1\u003c/sup\u003e. Compared to a best commercial graphite-based LIB (280 Wh kg\u003csup\u003e\u0026ndash;1\u003c/sup\u003e, 770 Wh L\u003csup\u003e\u0026ndash;1\u003c/sup\u003e), our AFLMBs eliminate the need for active anode materials, resulting in an 81% increase in gravimetric energy density, a 117% increase in volumetric energy density, and a 15-25% reduction in cost per kWh. Given that a battery life of 300 cycles effectively meets the operational requirements of drones, AFLMBs hold substantial promise for applications in drone and electric aircraft markets, where high gravimetric and volumetric energy and power densities are critical. With further advancements in fast charging and lifespan, AFLMBs could potentially be extended to the electric vehicle market in the future.\u003c/p\u003e\n\u003cp\u003e\u003cbr\u003e\u003c/p\u003e"},{"header":"Methods","content":"\u003cp\u003e\u003cstrong\u003eElectrolyte preparation\u003c/strong\u003e\u003cstrong\u003e\u0026nbsp;and their p\u003c/strong\u003e\u003cstrong\u003ehysic\u003c/strong\u003e\u003cstrong\u003eo\u003c/strong\u003e\u003cstrong\u003echemical properties\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eLiFSI, LiPF\u003csub\u003e6\u003c/sub\u003e, LiDFOB, LiBF\u003csub\u003e4\u003c/sub\u003e, DEC, FEC, FEMC, DME and commercial carbonate electrolyte ( 1.0 M LiPF\u003csub\u003e6\u003c/sub\u003e in EC/DMC ) were purchased from Duoduo Chem. HFE, TTE and NDFA was purchased from Tokyo Chemical Industry Co..\u0026nbsp;These solvents were dried by 3\u0026nbsp;\u0026Aring;\u0026nbsp;molecular sieves for two days\u0026nbsp;before use.\u0026nbsp;The\u0026nbsp;AFE electrolyte was prepared by dissolving 1.0\u0026thinsp;M LiPF\u003csub\u003e6\u003c/sub\u003e in FEC/FEMC/HFE\u0026nbsp;(weight ratio=2:6:2).\u0026nbsp;The LHCE electrolyte was prepared by dissolving 1.5\u0026thinsp;M LiFSI in\u0026nbsp;DME/TTE\u0026nbsp;(molar ratio=1.2:3). The\u0026nbsp;HCDE electrolyte was prepared by dissolving 2.0\u0026thinsp;M LiDFOB and 1.4 M\u0026nbsp;LiBF\u003csub\u003e4\u0026nbsp;\u003c/sub\u003ein\u0026nbsp;FEC/DEC\u0026nbsp;(volume ratio=1:2). The BAFF electrolyte was prepared by dissolving 1.6 M LiDFOB in\u0026nbsp;NDFA\u0026nbsp;(molar ratio=1:5).\u0026nbsp;Electrolytes were stored in an argon-filled glovebox (Mikrouna, oxygen \u0026lt;0.1 ppm, water \u0026lt;0.1 ppm) at room temperature.\u0026nbsp;The water content of the solution,\u0026nbsp;as measured by a Karl Fischer aquameter, was less\u0026nbsp;10\u0026thinsp;ppm.\u0026nbsp;The ionic conductivity was\u0026nbsp;measured\u0026nbsp;by AC impedance spectrometer\u0026nbsp;(Solartron,\u0026nbsp;1470E) in a symmetrical Pt|electrolyte|Pt cell.\u0026nbsp;The viscosity and density of\u0026nbsp;electrolytes\u0026nbsp;were measured using a kinematic viscometer (Anton Paar, SVM 3001).\u0026nbsp;The solution structure was studied by a\u0026nbsp;Raman spectrometer\u0026nbsp;(Anton Paar,\u0026nbsp;Cora\u0026nbsp;5700)\u0026nbsp;with an exciting laser of 785 nm.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAnode-free pouch cells construction\u003c/strong\u003e\u003cstrong\u003e\u0026nbsp;and e\u003c/strong\u003e\u003cstrong\u003electrochemical measurements\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eCu||NCM811\u0026nbsp;anode-free lithium pouch cell was assembled using double coated NCM811 (5.6 mAh cm\u003csup\u003e\u0026ndash;\u003c/sup\u003e\u003csup\u003e2\u003c/sup\u003e per one side) as the cathode,\u0026nbsp;commercial\u0026nbsp;polyethylene (PE)\u0026nbsp;as the separator,\u0026nbsp;and bare Cu foil without any surface coating as the anode.\u0026nbsp;All pouch\u0026nbsp;cells were filled with\u0026nbsp;1.5\u0026nbsp;g Ah\u003csup\u003e\u0026ndash;\u003c/sup\u003e\u003csup\u003e1\u003c/sup\u003e electrolyte and vacuum sealing in an argon glovebox.\u0026nbsp;Detailed cell parameters\u0026nbsp;are provided in\u003cstrong\u003e\u0026nbsp;Table S6\u003c/strong\u003e.\u0026nbsp;All anode-free pouch cells were conducted under galvanostatic charge\u0026ndash;discharge tests using NEWARE battery tester (NEWARE CT-4008-5V-6A). The constant-voltage charge process was applied until the charge current decayed to\u0026nbsp;0.05C.\u0026nbsp;The cycling\u0026nbsp;test was performed\u0026nbsp;at 30\u0026nbsp;℃\u0026nbsp;by\u0026nbsp;charging at 0.1C\u0026nbsp;and discharging\u0026nbsp;at 0.3C, after two formation cycles at\u0026nbsp;0.1C.\u0026nbsp;The temperature-dependent performance was tested by\u0026nbsp;charging at 0.1C,\u0026nbsp;30\u0026nbsp;℃\u0026nbsp;followed by\u0026nbsp;discharging at different temperatures (-40 ~ 60\u0026nbsp;℃).\u0026nbsp;The rate performance was\u0026nbsp;tested by\u0026nbsp;charging at 0.1C,\u0026nbsp;30\u0026nbsp;℃\u0026nbsp;followed by\u0026nbsp;discharging at different rates\u0026nbsp;using\u0026nbsp;a\u0026nbsp;NEWARE battery tester\u0026nbsp;(CT-8008-5V60A-NTFA).\u0026nbsp;A 1C rate corresponds to 5.6 mA\u0026nbsp;cm\u003csup\u003e\u0026ndash;\u003c/sup\u003e\u003csup\u003e2\u003c/sup\u003e. The resistances of\u0026nbsp;Li\u003csup\u003e+\u003c/sup\u003e transport through SEIs were measured by\u0026nbsp;a MPG-2 electrochemical workstation (Bio-Logic, France)\u0026nbsp;in a\u0026nbsp;symmetrical\u0026nbsp;Li||Li\u0026nbsp;cell. The Li electrodes were retracted from the charged Cu||NCM811\u0026nbsp;pouch cells after five cycles.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCharacterizations\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eAll the cycled samples were collected from pouch cells and rinsed by NDFA or DME solvent, followed by handling in transfer devices under argon atmosphere to avoid the air contamination. The morphologies of top, slope (30\u0026deg;) and bottom view of deposited Li samples were characterized using a field emission SEM (SU8230, Hitachi).\u0026nbsp;The cross-sectional morphologies\u0026nbsp;of deposited Li\u0026nbsp;samples\u0026nbsp;were\u0026nbsp;characterized using an\u0026nbsp;Cryo-FIB-SEM\u0026nbsp;(Helios 5 UX,\u0026nbsp;Thermo Fisher).\u0026nbsp;The operating voltage and emission current of the electron beam were 5 kV and 0.2 nA, respectively. A gallium-ion beam source (30 kV) was used to mill the sample with 2 nA for pattern milling, 26 pA for imaging, and 0.2 nA for cross-section cleaning. The stage temperature was maintained at \u0026minus;195 \u0026deg;C to prevent beam damage.The crystallographic information of deposited Li samples after five cycles was characterized by 2D GIXD with Eiger2 R 500k 2D detector (D8 Discover, Bruker) with an incident angle of 0.2\u0026deg;. The microstructure of the BAFF-derived SEI on Li surface (after five cycles) was characterized by cryo-TEM (JEM 2100F, JEOL). The\u0026nbsp;surface roughness and\u0026nbsp;Young\u0026rsquo;s modulus of SEIs\u0026nbsp;on Li surface were characterized by a\u0026nbsp;Cypher ES\u0026nbsp;instrument\u0026nbsp;(OXFORD)\u0026nbsp;using AM-FM module.\u0026nbsp;The probe was calibrated by\u0026nbsp;a standard sample of PVDF membrane with Young\u0026rsquo;s modulus of 2.45 GPa.\u0026nbsp;The chemical composition of BAFF-derived SEIs on bare Cu foils (after five cycles) were characterized by XPS,\u0026nbsp;MALDI-TOF-MS, and ss-NMR.\u0026nbsp;XPS experiments were conducted on a X-ray photoelectron spectrometer\u0026nbsp;with Al-K\u0026alpha; radiation\u0026nbsp;(Escalab 250Xi, Thermo Fisher).\u0026nbsp;A charge neutralizer was applied to compensate the sample surface charge.\u0026nbsp;MALDI-TOF-MS measurements were performed on\u0026nbsp;an\u0026nbsp;AXIMA-Performance\u0026nbsp;Shimadzu\u0026nbsp;Mass Spectrometer in linear mode (Power: 110 W) using 2,5-dihydroxybenzoic acid (DHB) as matrix.\u0026nbsp;Solid-state\u0026nbsp;\u003csup\u003e1\u003c/sup\u003eH,\u0026nbsp;\u003csup\u003e11\u003c/sup\u003eB\u0026nbsp;and\u0026nbsp;\u003csup\u003e19\u003c/sup\u003eF\u0026nbsp;Magic Angle Spinning (MAS)\u0026nbsp;NMR\u0026nbsp;experiments\u0026nbsp;were\u0026nbsp;performed\u0026nbsp;on a Bruker-AVANCE-500M\u0026nbsp;NMR\u0026nbsp;spectrometer\u0026nbsp;using a 1.3 mm double-resonance MAS probe at 60 kHz spinning rate. Dynamic-weighted analysis was applied to differentiate the polymer and small molecule components using the Carr-Purcell-Meiboom-Gill (CPMG) scheme for T\u003csub\u003e2\u003c/sub\u003e-filtering. The\u0026nbsp;amount of inactive metallic Li and LiH\u0026nbsp;in the fully deposited samples were quantified by MST experiments on a Hiden Analytical Mass Spectrometer (QIC-20) according to previously reported procedures\u003csup\u003e16\u003c/sup\u003e.\u0026nbsp;\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eData availability\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe data that support the findings of this study are available within this article and its Supplementary information. Additional data are available from the corresponding authors upon reasonable request.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAcknowledgements\u0026nbsp;\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThis work was supported by Research Center for Industries of the Future (RCIF), Zhejiang Key Laboratory of 3D Micro/nano Fabrication and Characterization, Westlake Education Foundation, and National Natural Science Foundation of China (Grant No. 21975207). The authors thank Drs. Yangjian Lin and Qike Jiang from Instrumentation and Service Center for Physical Sciences at Westlake University for supporting in Cryo-FIB-SEM and Cryo-TEM characterizations and data interpretation, and Profs. Pengfei Hu, Zexin Jin and Jiuan Lv at Westlake University for discussion on mechanistic understanding. Besides, the authors thank Drs. Xiaohe Miao, Lin Liu, Pei Sheng and\u0026nbsp;Ms. Huan Zhang (from Instrumentation and Service Center for Physical Sciences), Drs. Yinjuan Chen, Zhong Chen, Ms. Danyu Gu and Xin Li (from Instrumentation and Service Center for Molecular Science), Ms. Yingchun Wu (from Instrumentation and Service Center for 3D Micro/nano Fabrication) for their assistance in measurements at Instrumentation and Service Centers of Westlake University.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAuthor contributions\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eJ.W. and L.L. designed the experiments. L.L. prepared the materials, performed the electrochemical measurements, and carried out SEM-EDS, AFM, XPS, and Raman characterizations. Y.X. guided L.L. to carry out the MST measurements. X.L. and Y.X. carried out the NMR analysis. J.W. and L.L. prepared the manuscript. J.W. conceived and directed the project.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCompeting interests\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eJ.W. and L.L. are inventors of the published patents (CN202210110653.X, PCT/CN2022/138887).\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\n\u003cli\u003eArmand, M. \u0026amp; Tarascon, J. M. Building better batteries. \u003cem\u003eNature\u003c/em\u003e\u003cstrong\u003e451\u003c/strong\u003e, 652-657 (2008).\u003c/li\u003e\n\u003cli\u003eLin, D. C., Liu, Y. \u0026amp; Cui, Y. Reviving the lithium metal anode for high-energy batteries. \u003cem\u003eNat. Nanotechnol.\u003c/em\u003e\u003cstrong\u003e12\u003c/strong\u003e, 194-206 (2017).\u003c/li\u003e\n\u003cli\u003eSchmuch, R., Wagner, R., Horpel, G., Placke, T. \u0026amp; Winter, M. Performance and cost of materials for lithium-based rechargeable automotive batteries. \u003cem\u003eNat. 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However, AFLMBs face a long-standing challenge of short lifespan due to the harsh conditions of lacking excess lithium and an anode host. This issue is associated with uneven lithium deposition/dissolution, rooted in the micro-inhomogeneity and fragility of solid electrolyte interphase (SEI) on the lithium metal surface. Here, we present a practical 500 Wh kg–1-level AFLMB design with enhanced lifespan, achieved using an electrolyte of 1.6 M lithium difluoro(oxalate)borate in N,N-Dimethyltrifluoroacetamide solvent. The electrolyte-derived B-F-based polymer-rich SEI exhibits sub-nanometer homogeneity, high flexibility, and fast Li-ion conductivity, which spontaneously evolves a self-adaptive mesh-film structure that ensures uniform ion flux and large-volume-change accommodation, thereby realizing reversible planar lithium-orientated deposition/dissolution of 5.6 mAh cm–2. Consequently, a 2.7 Ah AFLMB (508 Wh kg–1, 1668 Wh L–1) without any host-material coating demonstrates stable cycling for 100 cycles at 100% depth of discharge (DoD) and 250 cycles at 80% DoD, with 80% capacity retention and an unprecedentedly high-power output of 2650 W kg–1 at 96 Wh kg–1. Our findings address the inherent structural instability of host-free electrodes, advancing the practical implementation of AFLMBs.","manuscriptTitle":"Planar lithium deposition/dissolution enabling practical 500 Wh kg–1 anode-free pouch cells","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-01-16 05:36:44","doi":"10.21203/rs.3.rs-5047161/v1","editorialEvents":[],"status":"published","journal":{"display":false,"email":"[email protected]","identity":"nature","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"nature","sideBox":"Learn more about [Nature](http://www.nature.com/nature/)","snPcode":"","submissionUrl":"","title":"Nature","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"ejp","reportingPortfolio":"Nature","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"2f75d4dc-2ee7-4f63-9cda-c523c36e365e","owner":[],"postedDate":"January 16th, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"published-in-journal","subjectAreas":[{"id":40594959,"name":"Physical sciences/Energy science and technology/Energy storage/Batteries"},{"id":40594960,"name":"Physical sciences/Materials science/Materials for energy and catalysis/Batteries"}],"tags":[],"updatedAt":"2026-04-23T07:16:03+00:00","versionOfRecord":{"articleIdentity":"rs-5047161","link":"https://doi.org/10.1038/s41586-026-10402-0","journal":{"identity":"nature","isVorOnly":false,"title":"Nature"},"publishedOn":"2026-03-17 04:00:00","publishedOnDateReadable":"March 17th, 2026"},"versionCreatedAt":"2025-01-16 05:36:44","video":"","vorDoi":"10.1038/s41586-026-10402-0","vorDoiUrl":"https://doi.org/10.1038/s41586-026-10402-0","workflowStages":[]},"version":"v1","identity":"rs-5047161","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-5047161","identity":"rs-5047161","version":["v1"]},"buildId":"8U1c8b4HqxoKbykW_rLl7","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}