# Stable cycling of high-voltage lithium metal batteries in ether electrolytes

## Abstract

The key to enabling long-term cycling stability of high-voltage lithium (Li) metal batteries is the development of functional electrolytes that are stable against both Li anodes and high-voltage (above 4 V versus Li/Li+) cathodes. Due to their limited oxidative stability ( <4 V), ethers have so far been excluded from being used in high-voltage batteries, in spite of their superior reductive stability against Li metal compared to conventional carbonate electrolytes. Here, we design a concentrated dual-salt/ether electrolyte that induces the formation of stable interfacial layers on both a high-voltage LiNi1/3Mn1/3Co1/3O2 cathode and the Li metal anode, thus realizing a capacity retention of >90% over 300 cycles and ~80% over 500 cycles with a charge cut-off voltage of 4.3 V. This study offers a promising approach to enable ether-based electrolytes for high-voltage Li metal battery applications.

## Main

With the fast-growing demands for high-energy storage, lithium (Li)-ion batteries (LIBs) can no longer satisfy the application needs due to their relatively low energy densities1,2. Nowadays, the majority of LIBs use a graphite anode coupled with a high-voltage (>4.0 V versus Li/Li+) Li+ intercalation/de-intercalation cathode. Replacing the graphite anode (372 mAh g−1 theoretical specific capacity) with a Li metal anode, which has a theoretical specific capacity of 3,862 mAh g−1 and a very low electrochemical potential (−3.040 V versus the standard hydrogen electrode), has been regarded as one of the most intriguing approaches to improve the battery energy density3,4,5,6. Nevertheless, the development of rechargeable Li metal batteries (LMBs) has been caught in a dilemma regarding the choice of electrolyte systems, which play an essential role in enabling the battery anode and cathode chemistries. Despite the formulation variations for different applications, the electrolyte solvents used in LIBs are predominately organic carbonate solvents (mixtures of cyclic and linear carbonates) because of their good stability on oxidative cathode surfaces and the effective solid electrolyte interface (SEI) layer on the graphite anode formed by ethylene carbonate7,8. Unfortunately, due to their poor cathodic stability, organic carbonate solvents are notorious for their incompatibility with Li metal anodes. Highly porous and dendritic Li deposition morphologies are commonly observed in organic carbonate electrolytes, which lead to very low Li Coulombic efficiencies (CEs) due to excessive side reactions and impose potential safety hazards due to dendrite penetrations9,10. On the other hand, among various other potential electrolytes for LMBs (including ionic liquids11,12, solid inorganic and polymer electrolytes13,14 and so on), ether solvents have always been regarded attractive because of their better reductive stability with Li metal, and thus higher Li CEs compared to carbonates15,16,17,18. However, the oxidation stability of ethers in regular 1 M salt concentration electrolytes is low ( <4.0 V), prohibiting their applications in practical high-voltage LMBs, especially with layer-structured cathode materials (LiMO2, M = Ni, Mn, Co) that have highly catalytic surfaces7,19. Previous tests of Li||LixMn2O4 using ether-containing electrolytes showed fast capacity fading and a very limited cycle life when charged to 4.3 V (ref. 20).

In this work, we demonstrate a highly efficient approach to achieve excellent cyclability of both a high-voltage cathode LiNi1/3Mn1/3Co1/3O2 (NMC) and the Li metal anode in an ether electrolyte by simply choosing suitable salts and electrolyte concentrations. Recently, various promising properties have been reported for concentrated electrolytes21. Significant progress has been made on suppressing dendrite formation in a variety of ether electrolytes and improving the CE of Li metal anodes. For example, CEs of >99.1% in an electrolyte of 4 M lithium bis(fluorosulfonyl)imide (LiFSI) in 1,2-dimethoxyethane (DME) have been reported by our group and others previously22,23. However, progress on stable ether electrolytes for high-voltage cathodes is very limited. In 2011, Watanabe and co-workers found that the electrochemical oxidation stability of triglyme and tetraglyme electrolytes with equimolar lithium bis(trifluoromethanesulfonyl)imide (LiTFSI) can be improved to ~5 V on the platinum (Pt) electrode24. Nevertheless, apparent capacity fadings in 200 cycles were still observed on a LiCoO2 cathode when charged to 4.2 V. Up to now, the demonstrations of LMBs in ether-based electrolytes were usually obtained with cathodes below 4.0 V (for example, LiFePO4)15,25. Here, we design functional electrolytes to enable stable cycling of state-of-the-art high-voltage NMC cathodes with catalytically active surfaces and avoid undesired ether oxidation under a high charge cut-off voltage of 4.3 V. At the same time, very efficient Li metal anode utilization is achieved. With the combination of these benefits, this study could greatly expand the possibilities for electrolyte design and promote the development of high-voltage LMBs with stable ether-based electrolytes.

## Cell performances with concentrated ether electrolytes

We first chose the concentrated LiTFSI–DME electrolyte with 3 M salt concentration (molar ratio LiTFSI:DME = 1:2, close to saturation, Supplementary Table 1) as a control electrolyte to study its stability on the NMC cathode (1.7 mAh cm−2 areal loading) with a charge cut-off voltage of 4.3 V. The viscosity and ionic conductivity of this electrolyte and other electrolytes studied in this work are listed in Supplementary Table 2. The battery can be cycled without over-charging even at 4.3 V cut-off voltage (Fig. 1a,b), indicating no excessive ether solvent oxidation happening in this concentrated electrolyte, which is consistent with the good anodic stability (~4.8 V on the Pt electrode, Fig. 1e) by the linear sweep voltammetry (LSV) scan. The significant improvement of the oxidation potentials can be attributed to two reasons. First, there are very limited free ether molecules in this concentrated electrolyte to contribute to the anodic current near 4.3 V. Second, the ether molecules bonded with Li+ cations may have a lower tendency to be oxidized because of the donation of the lone electron pairs on oxygen atoms24,26. Furthermore, the high LiTFSI concentration can inhibit the Al current collector corrosion (Supplementary Fig. 1a) because of the low solubility of the reaction product Al(TFSI)3 (ref.27). However, there is continuous capacity fading in the first 130 cycles before the fluctuation and fast decay of the capacity after 160 cycles, dropping to about 108 mAh g−1 after 250 cycles. When the cell was disassembled, it was found that the Li metal anode was covered by a thick surface layer. This surface layer was made of fine powders with high porosity and broke easily when peeling off the separator above (Supplementary Fig. 2). When the cycled Li metal anode was replaced with a fresh one and a new batch of electrolyte was added, the capacity of the cycled cathode in the new cell recovered to ~122 mAh g−1 (Fig. 1a), indicating the capacity fading of the original cell can be partially attributed to the decay of the Li metal anode because the average Li CE of the 3 M LiTFSI–DME is only 95.7% (Supplementary Table 1), measured in Li||Cu cells using a modified Aurbach’s method (see details in Methods). However, the limited capacity recovery and the continuous decay of the reassembled cell with a new Li anode also suggest the incompatibility of the NMC cathode with the 3 M LiTFSI–DME electrolyte, similar to the previous results24 on LiCoO2 cathodes at 4.2 V. Though the X-ray diffraction (XRD) pattern of the cycled cathode (Supplementary Fig. 3) shows no significant difference from the pristine cathode, an apparent cathode impedance increase was observed in the symmetric NMC||NMC cells using cathodes retrieved after 50 and 300 cycles (Supplementary Fig. 4a). It indicates the decay is mainly due to changes at the interface between the highly active cathode and the ether electrolyte.

To address the instability issue of ether-based electrolytes on the 4-V cathodes during charging, it is important to study components that can sacrificially passivate the cathode active sites and thus inhibit the catalytic oxidation of electrolytes. We chose lithium difluoro(oxalato)borate (LiDFOB) as such a passivation agent because of its ability to form an insoluble interface layer on high-voltage NMC cathodes in electrolytes, as well as its higher solubility than its analogue lithium bis(oxalato)borate (LiBOB)28. In the 4 M LiDFOB−DME electrolyte, the anodic current starts to increase from ~4.3 V and shows a peak before an exponential increase at ~5.2 V (Fig. 1e). The appearance of the current peak is probably due to the formation of a passivation film on the Pt electrode surface from the oxidation of LiDFOB starting around 4.3 V. Similarly, effective protection of the Al current collector can be realized with the LiDFOB oxidation, as observed in Supplementary Fig. 1b. Nevertheless, the cycling test of the Li||NMC cell using the concentrated 4 M LiDFOB–DME electrolyte shows only a relatively stable capacity during the first 20 cycles before a quick decay (Fig. 1a). The poor performance of the Li||NMC cell with 4 M LiDFOB–DME electrolyte can be mainly attributed to the instability of the Li anode (Li CE about 68.6%, Supplementary Table 1), and the continuous increase in cell overpotential and the capacity fading (shown in Fig. 1c) are mainly due to the accumulation of side products on the anode, which is supported by the Li replacement tests (Supplementary Fig. 5).

Significantly improved cycling performance was obtained when the 4 M dual salt–DME electrolyte (that is, 2 M LiTFSI + 2 M LiDFOB in DME) was employed, achieving a very high capacity retention of 90.5% (133 mAh g−1) after 300 cycles and 78.9% (116 mAh g−1) after 500 cycles with higher and more stable cell CE (Supplementary Fig. 6), though the average Li CE in this dual-salt electrolyte is only 94.6% (Supplementary Table 1). In addition, there is only a minimal increase of cell overpotential during the 500 cycles (Fig. 1d), which is mainly associated with the impedance increase of the Li metal anode, to be discussed below. Furthermore, the 4 M dual-salt–DME electrolyte shows better charging and discharging rate capabilities than the other two sole-salt electrolytes, as shown in Supplementary Fig. 7. It is demonstrated that the ether-based electrolyte can be cycled in Li||NMC batteries under the high charge cut-off voltage of 4.3 V.

## The cathode/electrolyte interface

X-ray photoelectron microscopy (XPS) depth profiling analysis on NMC cathodes was carried out to understand the reasons for the excellent cycling stability of the 4 M dual-salt–DME electrolyte over the other two electrolytes (4 M LiDFOB–DME and 3 M LiTFSI–DME) in Li||NMC cells. As shown in Fig. 2a, compared to the pristine NMC cathode, the C 1s spectrum of the NMC cathode after cycling in 3 M LiTFSI–DME shows little change. Common carbonaceous species, for example, C–C/C–H, C–O, C=O, as well as CH2–CF2 and CH2CF2 (from polyvinylidene fluoride (PVDF) binder), where the elements with underscores represent the atoms of interest in the identified species, are found in the pristine and the cycled NMC cathodes at different depths. The appearance of LiF in the F 1s spectrum suggests the decomposition of LiTFSI on the NMC surface under the high charging voltage. For the O 1s spectra (Fig. 2b), very similar compositions are found at different depths on the cathode cycled in the 3 M LiTFSI electrolyte (O = C–O, C=O and possible O–H), suggesting DME molecules are probably involved in the cathode electrolyte interface (CEI) layer formation process. In addition, an increase of the O 1s signal at ~530 eV is seen on depth profiling, similar to the situation on the pristine cathode. Therefore, this signal should come from the metal oxide (M–O) bonds, which indicates the 3 M LiTFSI–DME electrolyte is unable to effectively passivate the highly active cathode surface. Though most DME molecules solvate Li+ cations and are less prone to oxidative decomposition, the dynamic solvating/de-solvating processes may leave a small fraction of free DME molecules under the attack of a highly active cathode surface. The oxidation of these free ether molecules may generate undesired acidic species, causing continuous cathode capacity fading, as observed in Fig. 1a20.

In the 4 M LiDFOB–DME electrolyte, though no apparent film growth can be observed on the NMC particle surface after 50 cycles by scanning electron microscopy (SEM) (Supplementary Fig. 8c), the cathode surface shows apparent enrichments in O and B on depth profiling (a detailed comparison of atomic ratios is shown in Supplementary Table 3) compared to the cathode cycled in the 3 M LiTFSI–DME electrolyte, suggesting the decomposition of LiDFOB on the NMC surface. Accordingly, the C 1s spectrum shows an apparent enrichment in C=O (288.5 eV) species, and B–O, B–F and LiF species are found in the B 1s (Supplementary Fig. 9) and F 1s spectra (Supplementary Fig. 10). The M–O signal (O 1s spectrum) becomes more and more apparent after sputtering for more than 5 nm, suggesting the continuous decomposition of LiDFOB on the NMC cathode is sufficiently suppressed after a thin CEI layer is formed during the initial cycles. Similar to the function of the anode SEI layer, the CEI formed here can eliminate direct contact between the active cathode surface and ether molecules, and thus suppress the cathode side reactions. This is in agreement with our LSV result that LiDFOB could be oxidized at around 4.3 V and form an insoluble passivating film.

As for the 4 M LiTFSI–LiDFOB dual-salt concentrated electrolyte, the main XPS features of the surface film on the cycled NMC cathode are quite similar to those in the 4 M LiDFOB electrolyte, implying the major contribution of LiDFOB in forming the CEI layer. Similarly, the CEI layer formed in the 4 M dual-salt concentrated electrolyte is expected to be thin, as indicated by the gradual increase of the M–O signal on sputtering as well as the SEM image after cycling (Supplementary Fig. 8d). In addition, a weak signal at ~190.5 eV in the B 1s spectrum (Supplementary Fig. 9) is attributed to B–N species, which implies the participation of both LiTFSI and LiDFOB in the CEI layer formation.

More evidence of effective cathode protection in the 4 M dual-salt ether electrolyte can be obtained from the transmission electron microscope (TEM) characterization. As shown in Fig. 3a, the pristine NMC particle has a clear, well-defined layered structure with an interplanar distance indexed to the (003) crystal plane, with no additional surface layer observed. However, after 50 cycles in the three concentrated ether-based electrolytes, the CEI layers with different thicknesses can be clearly observed, and different electrolytes lead to different thicknesses of the CEI layers. The thicknesses of the CEI layers (average in our observed area) on the cycled NMC cathodes are about 23 nm for the 3 M LiTFSI–DME electrolyte (Fig. 3b), 6 nm for the 4 M LiDFOB–DME electrolyte (Fig. 3c), and about 4 nm for the 4 M dual-salt–DME electrolyte (Fig. 3d). The results indicate that the addition of LiDFOB enables efficient protection on the NMC cathode, especially for the dual-salt electrolyte, so further electrolyte decomposition is greatly suppressed during cycling. It is also seen from Fig. 3b that the CEI layer for the 3 M LiTFSI–DME electrolyte has a large domain of the rock-salt phase inside the CEI layer, which is quite different from the R$$\overline 3$$ m structure of the bulk NMC, though no apparent changes are found by SEM (Supplementary Fig. 8b). Though some regions of the CEI layer have a better crystallinity, the majority show a polycrystalline nature.

As mentioned earlier, the oxidation of ether molecules can generate acidic species at the cathode interface to induce transition metal ion dissolutions. Due to the high concentration of the electrolyte, these transition metal ions cannot be easily solvated and diffuse into the bulk electrolyte. Instead, they accumulate near the cathode surface and re-precipitate into this new interphase. In the 4 M LiDFOB–DME electrolyte, a much thinner and amorphous CEI layer was formed on the cathode surface, which is derived mainly from LiDFOB decomposition, as revealed by the XPS results, and could help inhibit further electrolyte corrosion to the cathode. However, a layer of ~7 nm thickness with a mainly rock-salt structure was found underneath the CEI layer. It may suggest there may be slow electrolyte penetration through the CEI layer to cause the gradual cathode decay. Interestingly, a partially crystalline CEI layer is found on the NMC surface when cycled in the 4 M dual-salt ether electrolyte. In addition, only a very thin layer (~4 nm) was transformed to the rock-salt phase, indicating an effective cathode protection can be achieved in the 4 M dual-salt electrolyte. The cathode impedance is also lower compared to others and shows no large changes after cycling (Supplementary Fig. 4). The above results indicate that there are some local/interfacial structural changes on the NMC particle surfaces during cycling, especially when cycled in the two sole-salt ether electrolytes. Importantly, the 4 M dual-salt ether electrolyte has caused very limited surface changes to the NMC cathode, demonstrating its efficient protection on the NMC cathode.

It is worth noting that though LiDFOB itself has the capability of forming a protective CEI layer on the NMC cathode surface, the salt concentration is also critical in maintaining the NMC cathode’s long-term stability. Lowering the total concentration of the dual-salt electrolyte would result in faster capacity decay (Supplementary Fig. 11) and increased cathode impedances (Supplementary Fig. 4b,c), which is probably related to cathode corrosion induced by side reactions of free DME molecules (Supplementary Fig. 12). The results prove that both electrolyte concentration and effective cathode passivation are significant in suppressing ether oxidation and maintaining stable operation of the highly active NMC cathode.

## The Li anode/electrolyte interface

In addition to the stability of an electrolyte on the cathode surface, another critical factor for the cell cycling stability is the electrolyte stability on the Li metal anode. SEM was used to characterize the morphologies of Li metal anodes after 50 cycles in the three electrolytes (Fig. 4). Even though the Li CE measured in Li||Cu cells using 3 M LiTFSI–DME is the highest among these electrolytes (Supplementary Table 1), the corrosion of the Li foil anode in Li||NMC cells was very serious (Fig. 4a), showing a corrosion layer more than 100 μm thick. Though the surface morphology looks like a dense layer that is formed on the Li anode (Fig. 4b), the anode surface layer is actually porous and is composed of numerous powders (Supplementary Fig. 2). It indicates that this surface layer cannot shield the bulk Li against electrolyte attack by the liquid electrolyte, resulting in continuous corrosion of the Li metal anode, regardless of whether it is a charge or discharge process. The Li metal anode in the 4 M LiDFOB–DME electrolyte also shows obvious corrosion with a surface layer thickness of about 120 μm (Fig. 4c). From the top-view SEM image (Fig. 4d), large pieces of isolated ‘dead’ Li are also observed, significantly different from that in 3 M LiTFSI–DME, which indicates the anode SEI layer formed in the concentrated LiDFOB electrolyte is highly non-uniform and explains the low Li CE measured as well as the quick cell capacity fading shown in Fig. 1a. Recently, Chen et al. verified the mechanism of the accumulation of dead Li and its direct effect on the capacity fade of the Li cells29. Importantly, the Li metal anode corrosion was greatly suppressed in the 4 M dual-salt electrolyte, where only the top layer (average about 22 µm) reacted after 50 cycles (Fig. 4e). The anode surface layer in the 4 M dual-salt electrolyte (Fig. 4f) is much more compactly integrated than that in the 3 M LiTFSI electrolyte, as well as more uniform than that in the 4 M LiDFOB electrolyte. The highly efficient Li metal cycling in the 4 M dual-salt ether electrolyte can be further demonstrated with a limited amount of Li metal (50-μm-thick Li on Cu as the anode). The Li||NMC cell with a 50-μm-thick Li anode can give more than 350 stable cycles in the 4 M dual-salt electrolyte, which is much better than those in the 4 M LiDFOB (~10 cycles) or 3 M LiTFSI electrolyte (~40 cycles) (Supplementary Figure 13). The Li metal anode CEs were calculated to be 98.4%, 59.1% and 91.7% for the 4 M dual-salt, 4 M LiDFOB and 3 M LiTFSI electrolytes, respectively, based on the initial Li mass and the accumulated capacity during cycling30. Based on the above observation, we found that though the average Li CE measurements in Li||Cu cells are convenient for electrolyte screening purposes, the Li plating/stripping behaviours and the short test time are different from those in Li metal full batteries. The long-term stability of the Li metal anode depends more on the ability of the anode SEI layer to accommodate Li volume fluctuations and to inhibit further electrolyte corrosion during repetitive cycling. Therefore, it is more realistic and important to test Li anode stability in actual Li metal batteries.

The mechanism of high Li CE in the 4 M dual-salt electrolyte can also be revealed by the XPS spectra of the cycled Li anodes (Fig. 5). In the 3 M LiTFSI–DME electrolyte, apparent reactions of Li metal with both LiTFSI and DME can be observed on the anode, as indicated by the formation of LiF (684.9 eV in F 1s) and the major presence of lithium alkyl oxides (R–OLi, 530.9 eV, O 1s). Apparent enrichments in the B element (B–F, B–O) and C=O species (C 1s) on the Li anode are observed in the 4 M LiDFOB–DME electrolyte, which is because the B–O bond in the DFOB anion has a bond strength (210.7 kJ mol−1) much weaker than other bonds, and the B–O bond strength in the DFOB radical is even at −50.1 kJ mol−1, as indicated by the theoretical calculations reported recently by our group31. Therefore, the B–O bond can break easily, especially in the radical case, and the SEI film contains the products from the reactions between LiDFOB and Li metal. Though the main features of the Li anode XPS spectra in the 4 M dual-salt electrolyte are close to those in the 4 M LiDFOB electrolyte, there are appreciable changes with the addition of LiTFSI. In the O 1s spectra, the 4 M dual-salt electrolyte renders a much lower amount of R–OLi species than the other two electrolytes, especially the 3 M LiTFSI electrolyte, while showing a larger ratio of components at ~533.1 eV, where both B–O and -(CH2–CH2–O)n- species are probably located6, which suggests the enrichment of -(CH2–CH2–O)n- species in the Li anode SEI layer. It is likely that the breakage of the B–O bond leaves the B atom back in the electron-deficient state, which has the tendency to coordinate with electron-rich species like other anions (for example, RO–Li from side reactions between DME and Li) to induce cross-linking or polymerization reactions32. These hypotheses can be validated by the Fourier-transform infrared (FTIR) spectra of SEI layers formed in different electrolytes, where shifts of C=O vibration signals in LiDFOB and prominent C–H signals are found in the 4 M dual-salt electrolyte (Supplementary Fig. 18 and the detailed discussion). Further density functional theory (DFT) calculations of the frontier orbitals of solvent, salts and their solvation complexes also support our discussions about the favourable electrolyte reactivities on both the anode and the cathode (Supplementary Fig. 19, Supplementary Table 4). Therefore, the formation of the polymeric SEI layer with the combination of LiDFOB and LiTFSI on anodes could improve the quality of the anode SEI layer for better Li metal protection.

It should be noted that the effective protection ability of the anode SEI layer formed in the 4 M dual-salt ether electrolyte has a potential trade-off in terms of efficient Li+ transport. With limited changes of the cathode structure (Supplementary Fig. 3) and impedance (Supplementary Fig. 4) observed, the latter gradual capacity decay is probably mainly from the polarization increases on the Li anode with repeated high-capacity high-rate Li plating/stripping6. A much better capacity retention was observed when using a lower cathode loading of 0.7 mAh cm−2 (Supplementary Fig. 20). Preliminary test results of the optimization of this dual-salt ether electrolyte show a much improved cell cycling performance (88% after 500 cycles) on increasing the LiTFSI/LiDFOB molar ratio to 7:3 and using a NMC442 cathode (same areal capacity of 1.7 mAh cm−2) with better stability under high voltages (Supplementary Fig. 21)33. Further optimization of this dual-salt ether electrolyte is underway to improve its compatibility with the Li anode.

## Conclusions

We have developed an ether-based electrolyte system for long-term cycling of Li||NMC batteries under charge cut-off voltages as high as 4.3 V, which breaks the long-standing voltage limitation for ether-based electrolytes. The ether-based concentrated electrolyte, 2 M LiTFSI and 2 M LiDFOB in DME, can greatly enhance the oxidation stability of ether molecules, effectively passivate the highly catalytic active NMC cathode surface under high voltage, as well as form a high-quality SEI layer on the Li metal anode for efficient Li metal cycling. All these fundamental findings extend the conventional knowledge on ether-based electrolyte systems, and provide an effective approach to achieve high-energy-density LMBs. It is important to realize that the practical implementation of LMBs requires high-loading cathodes with limited amounts of Li anode and electrolyte, which would incur further challenges to achieve a long cycle life of an LMB. Tests and further optimizations of this ether-based electrolyte aiming at practical applications are underway.

## Methods

### Materials

LiNi1/3Mn1/3Co1/3O2 (NMC) cathode material was purchased from BASF Battery Materials and used as received. The laminate of NMC electrode (~10.7 mg active material cm−2 or ~1.7 mAh cm−2) was prepared by casting a slurry mixture containing 96 wt% active material, 2 wt% C-NERGYTM SUPER C65 conductive carbon (Timcal), and 2 wt% polyvinylidene fluoride binder (Kureha L#1120) in N-methyl-2-pyrrolidone onto an aluminium (Al) current collector foil. NMC442 Cathode (LiNi0.4Mn0.4Co0.2O2) electrode laminates (active material 10.8 mg cm−2) were supplied by the Cell Analysis, Modelling, and Prototyping (CAMP) Facility located at Argonne National Laboratory (ANL). After drying, the electrodes were calendared and punched into disks with a diameter of 1.27 cm or an area of 1.27 cm2, and further dried at 75 °C under vacuum for 12 h. Battery-grade LiTFSI and DME were purchased from BASF as well. LiDFOB of battery grade was kindly provided from Materials Engineering Research Facility (MERF) at ANL. Lithium (Li) chips (250 µm thick, 1.56 cm diameter) and Al foil (15 µm thick) were purchased from MTI Corporation. Li foil of 50 µm thickness was acquired from Rockwood Lithium Inc. Copper (Cu) foil was ordered from All Foils, Inc. These chemicals were kept and handled in a glovebox (MBraun LABmaster) circulated with high-purity argon gas ( <1 ppm O2 and <1 ppm H2O).

### Electrochemical measurements

Charge/discharge performances were measured using CR2032 coin-type batteries. Due to the corrosion of stainless steel in 3 M LiTFSI–DME electrolyte, an Al-clad cathode case (MTI Corporation) was used for coin cell assembly. LMBs were constructed using a NMC positive electrode disc, a Li metal chip (250 μm thickness) as the anode, one piece of polyethylene (PE) separator (Celgard), and the prepared electrolyte (70 μl in each battery). All the Li||NMC batteries were tested on Land battery testers (Wuhan, China) within voltage windows of 2.7 and 4.3 V. LSV and CV studies of the electrolyte solutions were conducted in a three-electrode cell configuration using a CHI606E workstation.

A modified Aurbach’s protocol was used to test the average Li CE, as described in our previous articles30,34. A Li chip, a piece of PE separator and a Cu foil were sandwiched together in a CR2032 coin cell with 75 µl electrolyte. A Li film of 5 mAh cm−2 areal capacity was first electrodeposited onto the Cu foil and then fully stripped to a 1 V cut-off voltage. After another Li film (5 mAh cm−2) was deposited, then one-fifth of the Li capacity (that is, 1 mAh cm−2) was stripped and redeposited for 10 cycles. Finally, the Li film was fully stripped away. A current density of 0.5 mA cm−2 was used during this test.

### Characterizations

Morphological characterizations of cycled Li metal along with the energy dispersive X-ray spectroscopy (EDX) analysis were performed on a Helios focused ion beam SEM at 5.0 kV. For sample preparation, the cycled Li metal samples were soaked in pure DME for 1 h and then cleaned with fresh anhydrous DME several times to remove residual electrolytes before drying under vacuum. The cross-sections of Li metal were obtained by using a razor blade to cut the Li anodes. For TEM analysis, the NMC particles were sectioned by means of a Zeiss NVision 40 FIB/SEM dual beam system. The TEM specimens were prepared by the standard FIB lift-out procedure. After thinning by a 30 kV Ga ion beam, the specimens were further cleaned with the aid of a 5 kV Ga beam. Thus, most of the damage layer from 30 kV Ga ion beam would be removed. A JEOL 2100 F TEM was employed for the high-resolution imaging at a 200 kV accelerating voltage.

The XPS testing was carried out using a Physical Electronics Quantera scanning X-ray microprobe, which was fitted with a monochromatic Al Kα X-ray source (1,486.7 eV) for excitation. All samples were rinsed with anhydrous DME several times to remove residual electrolyte, then dried under vacuum. To avoid side reactions or electrode contamination with ambient oxygen and moisture, Li metal samples were transported from the glovebox to the SEM and XPS instruments in a hermetically sealed container filled with argon gas.

The FTIR spectra were measured by a Bruker Optics Vertex 70 FTIR spectrometer in diffusion reflectance mode (DRIFTS). The sample chamber was purged with N2 gas. The reference spectrum was from KBr powder. Each spectrum was averaged over 64 scans. The scan range was between 4,000 and 400 cm−1 with a resolution of 4 cm−1.

Electrochemical impedance spectroscopy (EIS) measurements were performed on a 1255B Solartron frequency response analyser and a 1287 electrochemical workstation in the frequency range from 100 kHz to 2 mHz, with a perturbation amplitude of ± 5 mV.

### DFT calculations

The chemical structures and the molecular orbital energies of the selected salt–solvent complexes were calculated by DFT using the generalized gradient approximation (GGA), as implemented in the Gaussian 09 suite of programs35. The B3LYP functional combined with the 6–311 + + G (d, p) basis set was used in geometry optimization calculations36,37.

### Data availability

The data that support the plots within this paper and other findings of this study are available from the corresponding authors upon reasonable request.

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## Acknowledgements

This work was supported by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the US Department of Energy (DOE) through the Advanced Battery Materials Research (BMR) program (Battery500 Consortium) under contract no. DE-AC02-05CH11231. The SEM, EDX, XRD, XPS and computational calculations were conducted in the William R. Wiley Environmental Molecular Sciences Laboratory (EMSL), a national scientific user facility sponsored by the DOE’s Office of Biological and Environmental Research and located at PNNL. PNNL is operated by Battelle for the DOE under Contract DE-AC05-76RLO1830. The LiDFOB salt was produced at the US DOE Materials Engineering Research Facility (MERF) and provided by K. Z. Pupek and T. L. Dzwiniel of ANL. The NMC442 electrode was made by the Cell Analysis, Modelling, and Prototyping (CAMP) Facility and provided by B. J. Polzin of ANL. The MERF and CAMP Facilities are fully supported by the DOE Vehicle Technologies Program within the core funding of the Applied Battery Research (ABR) for Transportation Program. This work was performed, in part, at the Center for Nanoscale Materials, a US Department of Energy Office of Science User Facility, and supported by the US Department of Energy, Office of Science, under contract no. DE-AC02-06CH11357.

## Author information

Authors

### Contributions

W.X. and J.-G.Z. proposed the research. W.X., S.J., X.R. and R.C. designed the experiments. S.J. and X.R. contributed equally to this work. S.J. and X.R. performed the electrochemical measurements with assistance from R.C. X.R., R.C. and J.Z. conducted the SEM and EDX observations. M.H.E. performed XPS measurements. Y.L. conducted TEM characterizations. D.H. performed IR measurements. W.Z. helped with XRD analysis. R.C. and J.-G.Z. participated in data analyses and discussion. D.M. and N.L. performed molecular dynamics calculations. Q.L. prepared the NMC electrodes. B.D.A. helped with the Li CE calculation with 50 μm Li foil. C.M. analysed the TEM results. S.J., X.R. and W.X. prepared this manuscript with inputs from all other co-authors.

### Corresponding authors

Correspondence to Ji-Guang Zhang or Wu Xu.

## Ethics declarations

### Competing interests

The authors declare no competing interests.

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## Supplementary information

### Supplementary Information

Supplementary Figures 1–21, Supplementary Tables 1–4, Supplementary References

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Jiao, S., Ren, X., Cao, R. et al. Stable cycling of high-voltage lithium metal batteries in ether electrolytes. Nat Energy 3, 739–746 (2018). https://doi.org/10.1038/s41560-018-0199-8

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