Ultrahigh Performance All Solid-State Lithium Sulfur Batteries: Salt

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Ultrahigh Performance All Solid-State Lithium Sulfur Batteries: Salt Anion's Chemistry-Induced Anomalous Synergistic Effect Gebrekidan Gebresilassie Eshetu, Xabier Judez, Chunmei Li, Maria Martinez-Ibañez, Ismael Gracia, Oleksandr Bondarchuk, Javier Carrasco, Lide M. Rodriguez-Martinez, Heng Zhang, and Michel Armand J. Am. Chem. Soc., Just Accepted Manuscript • DOI: 10.1021/jacs.8b04612 • Publication Date (Web): 15 Jul 2018 Downloaded from http://pubs.acs.org on July 15, 2018

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Ultrahigh performance all solid-state lithium sulfur batteries: salt anion's chemistry-induced anomalous synergistic effect Gebrekidan Gebresilassie Eshetu, Xabier Judez, Chunmei Li, Maria Martinez-Ibañez, Ismael Gracia, Oleksandr Bondarchuk, Javier Carrasco, Lide M. Rodriguez-Martinez, Heng Zhang,* and Michel Armand* Electrical energy storage department, CIC Energigune, Parque Tecnológico de Álava, Albert Einstein 48, 01510 Miñano, Álava, Spain. ABSTRACT: With a remarkably higher theoretical energy density compared to lithium-ion batteries (LIBs) and abundance of elemental sulfur, lithium sulfur (Li-S) batteries have emerged as one of the most promising alternatives among all the post LIB technologies. In particular, the coupling of solid polymer electrolytes (SPEs) with the cell chemistry of Li-S batteries enables a safe and high-capacity electrochemical energy storage system, due to the better process-ability and less flammability of SPEs compared to liquid electrolytes. However, the practical deployment of all solid-state Li-S batteries (ASSLSBs) containing SPEs is largely hindered by the low accessibility of active materials and side reactions of soluble polysulfide species, resulting in a poor specific capacity and cyclability. In the present work, an ultrahigh performance of ASSLSBs is obtained via an anomalous synergistic effect between (fluorosulfonyl)(trifluoromethanesulfonyl)imide anions inherited from the design of lithium salts in SPEs and the polysulfide species formed during the cycling. The corresponding Li-S cells deliver high specific/areal capacity (1394 mAh gsulfur–1, 1.2 mAh cm−2), good coulombic efficiency, and superior rate capability (~800 mAh gsulfur–1 after 60 cycles). These results imply the importance of the molecular structure of lithium salts in ASSLSBs and pave a way for future development of safe and cost-effective Li-S batteries.

1. INTRODUCTION In today’s modern and energy-conscious society, the quest of energy storage for using in large-scale applications such as in the electro mobility (xEVs, EVs/HEVs/PHEVs), and renewable or dispersed energy storage has become compulsory.1,2 Amid existing large spectrum energy storage devices, lithium-ion batteries (LIBs) have dominated as key-enabling technologies in the market. However, the existing practical LIBs fall far below the stringent requirements of the above-mentioned demands where a high energy density, beyond the current capability of LIBs, is needed. Historically, in comparison to computer industry, which enjoyed a doubling in memory capacity every 18 months as indicated by Moore's law, the rate of progress in energy storage has been quite lethargic.1,2 In the past 150 years, the practical energy density of commercial batteries has increased only by 6-fold, from the first-generation lead-acid batteries (∼40 Wh kg–1) to the contemporary LIBs (∼240 Wh·kg–1).3 Thus, though the evolutionary progress has increased the knowledge of acquisition and the degree of the battery technology´s maturity significantly, a new paradigm shift towards the development of new energy storage devices is urgently needed. With its exceptional high specific capacity (3860 mAh g–1), lowest electrochemical potential (−3.04 V vs. standard

hydrogen electrode, SHE), lowest density (0.53 g cm−3) and lightest nature (6.94 g mol−1), lithium metal (Li°) is regarded as a “holy grail” electrode and accordingly has received extensive attention from both the academic and industry communities.4−8 Among all the Li°-based rechargeable batteries (LMBs), lithium-sulfur (Li-S) technologies are considered to be the most auspicious candidates for the next-generation energy storage systems.9−13 This is attributed to their overwhelming benefits such as high theoretical energy density (2600 Wh kg−1) computed on the basis of Li° anode and cyclo-octa sulfur (S8) cathode, abundance of sulfur resources, low cost, and environmental benignancy.9−13 However, the practical deployment, in light of the above beneficial features, is hampered by several intrinsic problems resulting from the complex Li-S cell chemistry, enlisting Li dendrite/mossy growth and possible cell short circuit, electronically insulating nature of S8 (ca. 10−30 S m−1) and of its lower order reduction products (e.g., Li2S, ca. 10−14 S m−1), soluble long chain polysulfides (PS) and their shuttling between cathode and anode, large volume expansion from S8 to Li2S (ca. 80%) etc.9−13 In order to address the above-mentioned dilemmas and improve the safety of Li-S batteries, considerable amount of work has been devoted to all solid-state Li-S batteries (ASSLSBs).14 The use of solid polymer electrolytes (SPEs) surmounts the safety-induced hazards (enlisting both

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thermal and chemical threats) arising from the highly flammable/combustible organic ether-based solvents employed in conventional liquid electrolytes.15 The success of SPEs has been well demonstrated by their preliminary application as electrolytes of solid-state rechargeable LMBs for an electric car, Autolib®.16 Moreover, the low density of SPEs, compared to other solid electrolytes like garnets, allows a high gravimetric energy density to be readily achieved. Notwithstanding, the practical energy density and cyclability of SPEs-based ASSLSBs are still far from the expectations.14,17 One of the strategies to overcome such inter-linking challenges of ASSLSBs lies in the in-depth understanding of the salt anion chemistry and thereby a proper selection of lithium salts for gaining robust SPEs.18,19 The anion chemistry dictates the nature and quality of both Li° anode and S cathode passivation layers, possible interaction/reaction with PS, wettability of electrolytes (i.e., degree of sulfur utilization), and so on. Lithium bis(trifluoromethanesulfonyl)imide (Li[N(SO2CF3)2], LiTFSI) is among the most widely used salts in the electrolytes for Li-S batteries; however, attempts to use LiTFSI in SPE-based ASSLSBs are grappled with the infinite charging, associated to PS shuttling even in the earliest cycles, though the first discharge capacity is as high as 900 mAh gsulfur–1.20 This is due to the inferior quality of the solid electrolyte interphase (SEI) layer formed on Li° anode.20 Lithium bis(fluorosulfonyl)imide (Li[N(SO2F)2], (LiFSI)), an analogue of LiTFSI, confers improved compatibility with Li° anode which is attributed to the formation of LiF-rich SEI layer, thus resulting in a stable cycling performance. Whereas, LiFSI-based cells have lower initial discharge and areal capacities, ascribed to the poor molten-state wettability of ‒SO2F compared to ‒SO2CF3 moiety in LiTFSI and/or possibly due to an irreversible reaction with PS.20,21 This shows that both can hardly be used as single electrolyte salts in ASSLSBs. Thus, a synergistic effect combining the beneficial features of both anions (i.e., FSI− and TFSI−) could be of paramount importance to build ASSLSBs with robust SEI layer, improved discharge/areal capacity, stable long-term cyclability, high coulombic/energy efficiency, etc. In this work, lithium (fluorosulfonyl)(trifluoromethanesulfonyl)imide (Li[N(SO2CF3)(SO2F)], LiFTFSI), having both ‒SO2CF3 and ‒SO2F functionalities (Figure 1), is proposed as an elegant salt anion to merge the complementary advantages of the TFSI− and FSI− in ASSLSBs. Unprecedentedly, Li-S cells using LiFTFSI/poly(ethylene oxide) (PEO) electrolyte can deliver an extremely high specific discharge capacity of 1394 mAh gsulfur–1 (83.2 % of the theoretical capacity), a high areal capacity of 1.2 mAh cm–2 with good coulombic efficiency, and superior rate capability. The possible mechanism of such a superior enhancement in the cycling performance of LiFTFSI-based ASSLSBs is investigated with the help of experimental and computational studies.

Figure 1. Comparison of the chemical structures for the LiFTFSI, LiFSI, and LiTFSI salts.

2. EXPERIMENTAL SECTION Preparation of polymer electrolyte membrane. The electrolyte membranes with an average thickness of 50 μm were prepared by a typical solvent casting method, followed by hot-pressing (High Temperature Film Maker Controller, Specac). Briefly, a pre-weighed amount of PEO was dissolved into acetonitrile and then the corresponding battery grade lithium salt (such as LiFTFSI (Provisco, Czech Republic), LiTFSI (Solvionic, France), LiFSI (Suzhou Fluolyte, China)) was added. In all the prepared electrolytes, the salt concentration was fixed at the optimized molar ratio of −CH2CH2O− (EO)/Li+ = 20:1. X-ray diffraction. Powder X-ray diffraction (PXRD) patterns of the prepared polymer electrolytes were recorded on a Bruker D8 Discover X-ray diffractometer, using λCu-Ka = 1.54056 Å radiation in the 2θ range from 2° to 80° with a step width of 0.0198°. Thermal behavior. The phase transitions of the polymer electrolytes were measured on a differential scanning calorimeter (DSC) (Q2000, TA Instruments). A protocol containing two consecutive scans at a cooling and heating rate of 10 oC min−1 in the temperature range of −80 to 100 o C was used. At the two temperature extremes, low (−80 o C) and high (100 oC) ends of each scan, the sample was allowed to stand for five minutes. Thermogravimetric analysis (TGA) was performed on a STA 449 F3 system connected to QMS 403 Aëolos (Netzsch). The samples were heated from room temperature to 600 oC at a heating rate of 10 oC min−1 under Argon flow. Ionic conductivity. The ionic conductivity of polymer electrolytes was measured by AC impedance spectroscopy using a VMP3 potentiostat (Biologic). The frequency ranged from 10−1 to 106 Hz with signal amplitude of 10 mV. A CR2032 type coin cell using two stainless steel (SS) blocking electrodes (SS | SPEs | SS) was assembled in an Argon filled glovebox and used for the measurement. Anodic stability: The experiments of the linear sweep voltammogram (LSV) of the polymer electrolytes were performed with a VMP3 potentiostat (Biologic) using a two-electrode cell at 70 oC. Stainless steel (surface area:


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0.07 cm2) served as working electrode, Li disk was used as both counter and reference electrodes. The LSV measurements were carried out between the open circuit potential (OCP) and 6.0 V vs. Li/Li+ at a scan rate of 1 mV s−1. Electrochemical stability of electrolyte/Li° electrode. Li° symmetrical coin cells (Li° || Li°) were assembled in an argon-filled glove box to investigate the electrochemical stability of electrolyte/Li° interphase. The galvanostatic cycling of the Li° symmetric cells was evaluated using a Maccor Battery Tester (Series 4000). The Li° symmetric cells were cycled galvanostatically at a current density of 0.1 mA cm−2, wherein the duration of each halfcycle was 2 h. Surface morphology and composition of Li° deposits. Surface morphologies of the Li° deposits were examined by a field emission Quanta 200 FEG (FEI), operated at 20 kV. The Li° deposits were obtained by the galvanostatic deposition of Li° on Cu substrates using Li | LiX/DME (X = FTFSI, FSI and TFSI) | Cu cells at a current density of 0.1 mA cm−2 for 20 hours. The compositions of the surface layer were measured by a Phoibos 150 X-ray photoelectron spectroscopy (XPS) with a nonmonochromatic Mg Kα source (hν = 1253.6 eV). The spectra were recorded with high resolution scans at low power (100 W, 20 eV pass energy, and 0.1 eV energy step). The calibration of the binding energy was performed taking into account as reference the graphitic signal at 284.4 eV. The samples were gently rinsed with 1,2-dimethoxyethane (DME) and dried thoroughly under vacuum before transferred to the SEM or XPS chamber by a home-designed airtight setup. Electrochemical reduction simulation. Biphenyl was dissolved in ultra-pure and dry tetrahydrofuran (THF) for about 40 minutes. Following, grounded Li° was added to the biphenyl/THF solution in a 1:1 mole ratio (biphenyl:Li). The reaction mixture was stirred for about 4 hours at room temperature using a special glass-coated magnetic stirring bar. Afterwards, a dark blue color, characteristic of the biphenyl radical anion was observed. For the two electrons extraction, biphenyl: Li of 1:2 by molar was used, resulting in the formation of a deep dark blue biphenyl radical di-anion. Reaction of polysulfide with anions. The solution of 0.1 M Li2S6/DME had been prepared by mixing Li2S and S with a mole ratio of 1:5 in DME, followed by stirring for one week at room temperature.22 The reaction between polysulfide species and the three investigated anions were carried out in DME solution at room temperature. To a stirred solution of 1 M LiX/DME (X = FTFSI, FSI and TFSI), a predetermined amount of 0.1 M Li2S6/DME solution with a molar ratio of LiX/Li2S6 of 200 was added. The UVVis measurements were carried out, without dilution, on a Cary 5000 UV-Vis spectrophotometer (Varian). DFT calculations using the Becke’s three parameters (B3) exchange functional along with the Lee-Yang-Parr (LYP) non-local correlation functional (B3LYP),23,24 as implemented in the numeric atom-centered basis set allelectron code FHI-aims.25,26 The starting geometries of the investigated PSAs considering different conformers, in

general, by varying either S‒N‒S‒F or S‒N‒S‒C dihedral angles. The initial geometries were constructed using the open-source molecular editor and visualizer Avogadro.27 The universal force field and genetic algorithm search tool as implemented in Avogadro to pre-screen dozens of lowenergy geometries was used. Then, their atomic structures using a trust radius method enhanced version of the Broyden-Fletcher-Goldfarb-Shanno optimization algorithm28 were fully relaxed. The “tight” settings, including the “tier2” standard basis set in the FHI-aims code for Li, C, N, O, F, and S atoms, were used, and the following thresholds for the convergence criteria were sett: 0.01 eV Å−1 for the maximum residual force component per atom in all structural relaxations, 10−4 electrons for the electron density, and 10−6 eV for the total energy of the system. S cathode preparation and cycling of Li-S polymer cell. Composite sulfur cathode was prepared with 40 wt% elemental sulfur (99.5 wt%, Sigma-Aldrich), 15 wt% conductive carbons (Ketjen Black, KJ600, Akzo-Nobel), and 45 wt% LiX/PEO (X = FTFSI, FSI and TFSI) as electrolyte. The S loading was from 0.9–1.1 mg cm–2. The procedure has been detailed in our previous work.20,29 Li-S polymer cells were assembled in an Argon filled glovebox using the prepared electrode as cathode, polymer electrolyte membrane as both electrolyte and separator, and Li metal disk (China energy Lithium) as anode. The cells were then cycled galvanostatically at a constant current (CC) mode between 1.6 and 2.8 V at 70 °C using a Maccor Battery Tester (Series 4000).

3. RESULTS AND DISCUSSIONS Physicochemical and electrochemical properties of SPEs. Figure 2 displays the fundamental physicochemical and electrochemical properties of LiX/PEO (X = FTFSI, FSI and TFSI) electrolytes at the molar ratio of EO/Li = 20. Thanks to the high molecular weight of PEO matrix, a free-standing and translucent membrane with a thickness of 50 μm could be obtained for the SPEs using the three salts, as exemplified by the optical appearance of LiFTFSI/PEO in Figure 2a. With the increase of salt content from EO/Li = 64 to 16, the LiFTFSI/PEO membranes become soft and sticky and loss their mechanical strength (Figure S1). The XRD-patterns (Figure 2b) of the three electrolytes show only two characteristic peaks, which belong to the crystalline phase of PEO (2θ =19.3 and 23.7), without any peaks that could be assigned to the correspondent lithium salt. This indicates the semi-crystalline nature of the PEO matrix, as well as a full solvation of lithium salt in the amorphous region of PEO via the complexion of Li+ cation with the strong electron-donating ethylene oxide units (donor number of 22) and a favorable entropy factor. The obvious endothermic peaks in the DSC traces (Figure 2c) at ca. 65 oC, resulting from the melting transitions of PEO, clearly confirm the existence of crystalline phases in those electrolytes. In addition, both LiFTFSI/PEO and LiTFSI/PEO electrolytes have lower crystallinities (χc) compared to LiFSI/PEO electrolyte (e.g., χc = 54%, 42%, and 59% for LiFTFSI, LiTFSI, and LiFSI/PEO, respectively),



due to the sluggish kinetics of crystallization in the presence of ‒SO2CF3 moieties in both FTFSI− and TFSI−.21 As one of the most important features, thermal stability of SPEs has a pivotal effect on the safety, durability and long-term stability of ASSLSBs. Individually, the three salts and PEO matrix have decomposition temperatures higher than 200 oC, thus yielding a good thermal tolerance of the corresponding SPEs (Figure 2d). LiFTFSI/PEO with a decomposition temperature at 220 oC (5 wt% mass loss) could sufficiently meet the requirement for its application in LMBs.

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The temperature dependence of ionic conductivities for the three SPEs shows a notable change at 50−60 oC (Figure 2e), due to the melting of crystallized PEO phases, which is supported by the DSC results where melting transitions around 65 oC are observed. At an operational temperature for Li-S cells, the ionic conductivity of LiFTFSI/PEO electrolyte maintains very close to that of LiTFSI/PEO (e.g., 6.5 × 10−4 S cm−1 (LiFTFSI/PEO) vs. 7.0 × 10−4 S cm−1 (LiTFSI/PEO) at 70 oC).

Figure 2. Physicochemical and electrochemical properties of LiX/PEO (X = FTFSI, FSI and TFSI) electrolytes. (a) Optical image of LiFTFSI/PEO membrane. (b) XRD patterns and (c) DSC traces of the three electrolytes. (d) TGA traces for the three electrolytes (top), and their neat components (bottom), (e) Arrhenius plots of ionic conductivity and (f) anodic stabilities at 70 oC for the three electrolytes. The salt content effect on phase transition and ionic conductivity of the LiFTFSI/PEO electrolytes is systematically studied. The melting transition of PEO turns to be lowered with the addition of LiFTFSI (Figure S2) due to the plasticizing effect of the sulfonamide salts, whereas

the optimal ionic conductivity is obtained with EO/Li = 20 and 12 (Figure S3) due to the trade-off between ion mobility and the number of charge carriers. The sample with EO/Li = 20 is a compromise between the ionic conductivity and mechanical property.


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The anodic stabilities of the three SPEs are presented in Figure 2f. The linear sweeping voltammetry profiles show relatively low anodic currents of < 10 μA cm−2 at 3.8 V vs. Li/Li+, which can be mainly ascribed to the oxidation of PEO matrix. Then, pronounced increase of currents, generated due to oxidation of the salt anions, are observed once the potentials exceed 5 V vs. Li/Li+. In principle, those anions are electrochemically stable and can be coupled with 4 V class cathode materials (such as lithium nickel manganese cobalt oxide (NMC), lithium nickel cobalt aluminum oxide (NCA)); the electrochemical stability of the SPEs is believed to arise from the decomposition of PEO matrix. Nevertheless, those PEO-based SPEs are ideal candidates for cathodes working at a potential lower than 3.8 V, such as lithium iron phosphate (LFP) and S cathodes. Among tested electrolytes, LiFTFSI/PEO shows a good stability up to ca. 5.3 V vs. Li/Li+ with minimal currents of < 20 μA cm−2, though the electrochemical decom-

position of FTFSI− becomes more prominent at 5.5 V vs. Li/Li+. With LiFTFSI, the residual current related to PEO oxidation is suppressed, and this could be attributed to the formation of a LiFTFSI oxidation-derived passivation layer on the cathode, which then kinetically hinders PEO oxidation. Interfacial stability with Li° electrode. The operation of Li° electrode is notoriously known for the formation of dendritic/mossy lithium which could lead to a decreased coulombic efficiency, poor cyclability and safety-induced concerns. Hence, a good interfacial stability between Li° electrode and electrolytes is crucial for achieving highperformance and safe ASSLSBs.9−14 Figure 3a presents the response of voltage vs. time for the Li° | LiX/PEO | Li° symmetrical cells upon continuous galvanostatic cycles. For those three cells with different lithium salts, stable voltage profiles are observed in the first few cycles, due to the formation of SEI layer on the surface of Li° electrode.

Figure 3. Electrochemical behavior of Li° electrode in the as-prepared electrolytes. Galvanostatic cycling of Li° symmetric cells at 0.1 mA cm−2 (half cycle time 2 h) using (a) polymer electrolytes at 70 oC and (b) liquid electrolytes at 25 oC. (c) Galvanostatic polarization of Li symmetric cells at 0.1 mA cm−2 for the liquid electrolytes at 25 oC. Both the LiFTFSI- and LiFSI-based Li° symmetric cells demonstrated very stable evolutions of voltage without any erratic values up to ca. 200 hours, while LiTFSI-based one fails after being cycled only for 50 hours. The same tendency is verified by the galvanostatic cycling of Li° symmetric liquid-based cells, as shown in Figure 3b, where 1,2-dimethoxyethane (DME), a low molecular weight version of PEO, is used as solvent for dissolving the lithium salts and thereby mimic the evolution in the corresponding LiX/PEO electrolytes. The liquid Li° sym-

metric cells using LiFTFSI and LiFSI salts can be cycled for more than 1400 hours, which is 30 times higher than LiTFSI-based liquid cell. These clearly suggest that the FSO2‒ functional group can greatly regulate the morphology, composition and mechanical stability of the SEI layers in both liquid and polymer recipes, thus mitigating the growth of Li° dendrites in the corresponding electrolytes. Interestingly, one may note that LiFTFSI-based Li° symmetric cells in both liquid and polymer electrolytes possess the lowest value of voltage (i.e. over potential) after



the first few stabilization cycles, e.g., LiFTFSI/PEO (28 mV) < LiFSI/PEO (48 mV) < LiTFSI/PEO (62 mV), which implies that the SEI layer formed on Li° electrode in the LiFTFSI- based electrolytes is less resistive compared to the ones formed in the other two salts-based electrolytes. This can be rationalized by the slightly lower tendency in electrochemical reduction for FTFSI‒ vs. FSI‒ and higher liability compared to TFSI‒, thus resulting in a passivation layer with an optimized LiF amount (see XPS analysis and chemical simulation in Figures. 5 and 6 for details). Figure 3c shows the galvanostatic Li° deposition measured in the Li° | LiX/DME | Li° cells. The LiFTFSI-based cell again exhibits superior stability during the Li deposition process, a high areal capacity of > 75 mAh cm−2 could be reached without any oscillation in voltage. This result further suggests that the SEI layer formed in the LiFTFSIbased electrolyte has a robust quality and allows the continuous plating of Li+ ions on Li° electrode. To shed some lights on the role of salt anions in dictating the electrochemical compatibility between Li° electrode and electrolytes, post-mortem analyses of the cycled Li-Cu cells were carried out by means of scanning electron microscopy (SEM) and X-ray photoelectron spectroscopy (XPS). DME-based liquid electrolytes are selected for simulating the PEO-based polymer electrolytes, since the cycled electrodes could be hardly recovered without any damage in the latter cases due to the strong

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adhesive properties of PEO-based SPEs towards Li° electrode. This is also noticed in a recent work on the PEObased SPEs containing lithium azide as an electrolyte additive.29 Figure 4 shows the SEM images of the Li° deposits on the surface of Cu electrodes obtained from those three liquid electrolytes. For LiFTFSI/DME and LiFSI/DME electrolytes, flake-like Li° deposits with a relatively homogenous distribution on Cu substrates are observed, though they are accompanied by a certain amount of needle-like deposits, possibly Li° dendrites. In contrast, large amount of dendritic Li° deposits are unevenly plated on the surface of Cu electrode with LiTFSI-based electrolyte, as observed in previous works.29,30 The morphological observation evidences that LiFTFSI and LiFSI salts are favorable for forming less dendritic, more dense and compact Li° deposits, thus resulting in better cycling performance in Li° symmetric cells (Figure 3). The different lithium plating/stripping performances and SEM morphologies in the different electrolyte systems could be ascribed to the nature of the SEI layer formed on the surface of Li° electrodes. To ratify our surmise, XPS analysis of Li° electrode surfaces was carried out. Based on the work from Aurbach group and our recent study,29,31 it could be concluded that DME and the above-mentioned salts can be reduced, though their liability towards reduction significantly varies. Figure 5 shows the Li1s, C1s, F1s spectra harvested from the outermost surface of the SEI layer formed in the three electrolyte systems.

Figure 4. SEM images of Li deposited onto Cu substrates at 0.1 mA cm‒2 (plating time 20 h) in LiX/DME (X = FTFSI (a, d), FSI (b, d) and TFSI (c, f)) electrolytes at 25 oC.


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The assignments of each peak in the XPS spectra as extracted from literature are summarized in Table S1, and the XPS survey spectra are also presented in Figure S4. The Li1s binding energy peak at around 57 eV is characteristic of the different Li-containing SEI-building materials such as LiF and could not be that of metallic Li as the binding energy for Li° is found at a lower value (ca. 52.3 eV).32 Armand et al. proposed a two-electron reduction mechanism of DME leading to the formation of H2C=HC‒O‒CH3, CH3OLi and LiH (Scheme 1 d-1).29 Aurbach and co-workers suggested a two-step two electron reduction

mechanism of DME with CH3OLi and H2C=CH2 as SEI building species.31 According to the authors, the reduction involves first the formation of a Li+ cation stabilized radical anion, which then decomposes to alkoxide, ROLi (CH3OLi) and methoxy ethyl radical. The later in turn undergoes a one electron transfer resulting in additional CH3OLi and ethylene (C2H4) (Scheme d-2). Hu et al.22 reinforced further the proof for the formation of RO‒Li species and oligomers with ‒OLi end group in DME-based electrolytes on Li° electrode. One of the common features for all tested salts in the C1s

Scheme 1. The proposed mechanism for the reduction of the electrolyte components (i.e., salts and solvent).



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photoelectron spectra is the presence of a signal at ~286.6 eV, attributed to carbon atoms surrounded by one oxygen atom (i.e., C‒O), such as CH3OLi generated from the electrochemical reduction of DME (scheme d-1 and d-2).

Figure 5. (a) XPS spectra of Li° deposited onto Cu substrates at 0.1 mA cm‒2 (plating time 20 h) in LiX/DME (X = FTFSI, FSI and TFSI) electrolytes at 25 oC. (b) Atomic concentration of Li, S, F and N on the Li° deposits as a function of etching time using the three liquid electrolytes. In the figure, R1C‒O and R2C‒O refer to H2C=HC‒O‒CH3 and CH3O‒ respectively. Thus, the spectra confirm the existence of DME-derived contribution to the passivation layer. In the case of LiTFSI and LiFTFSI electrolytes, they presented an additional

broader signal in range of 287.0 to 287.7 eV, which could be ascribed to the presence of R1C‒O (H2C=HC‒O‒CH3). Considering the intensity of this peak, the SEI layer


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formed in LiTFSI/DME can be regarded as mainly dominated by solvent (i.e. DME) reduction products. Analogous to the organic building species of the SEI layer, the composition of the inorganic layer largely depends on the nature of the electrolyte composition, mainly emanating from the decomposition of the electrolyte salt and/or residual salts trapped in the passivation layer. The F1s signal at 686.5 eV corresponds to LiF and its formation is favored in the order of LiFSI > LiFTFSI > LiTFSI (see Table S2 for the comparison of normalized percentage of each F-containing specie). That is, LiF is the predominant SEI species in LiFSI, and only trace amount could be detected on the SEI layer originating from LiTFSI-based electrolyte. As reported in literature,32 the differences are ascribed to the bond strength between S‒F and C‒F. The S‒F bond in FSI‒ is more labile towards reduction compared to C‒F in TFSI‒ or to one side of FTFSI‒, and accordingly, LiF is the dominant species. LiFTFSI/DME presents an intermediate tendency in forming LiF, which is in line with its structure where it is endowed half with C‒F and half with S‒F bond functionality. The electrochemical reduction mechanisms of LiFSI, LiFTFSI and LiTFSI are depicted in Scheme 1. The high and relatively midway amount of LiF observed for LiFSI and LiFTFSI respectively is further verified by the signal at higher binding energy, ca. 690.0 eV, which corresponds either to the trapped pristine salts and/or their incomplete decomposition products. Aiming at substantiating the above Li1s, C1s and F1s spectral observations and accompanying claims, depthdependent elemental compositions are given in Figure 5b, and the corresponding XPS survey spectra are also presented in Figure S5.

The Li signal depicts an increased concentration value for all the three electrolytes with increasing the sputtering depth, due to the removal of the outermost reduction products from salts and solvents. As shown in Figure S6, highly reductive species such as LiF tend to be the pronounced components on Li electrode at the inner part of the SEI layers. Moreover, the atomic concentration of Li is found to be higher than that of fluorine (F1s) implying that most of the lithium is neither bound to sulfur (S2p) nor to nitrogen (N1s) containing SEI species and this upholds well with the low concentrations of the respective spectra. This once again testifies that the passivation layer in the case of imide salts having S‒F bond is dominated by LiF, as a principal building material. In summary, one can assume that the SEI layer in LiFTFSI-based electrolyte contains an optimum amount of inorganic species (mainly LiF), a prerequisite for mechanically and electrochemically stable SEI layer. This could explain the improved stripping-plating and most importantly the less resistive behavior of the passivation layer. Mechanistic understanding of the anions chemistry on Li° electrode. Targeting at simulating the above claimed electrochemical reductions of the salts on Li° surface and thereby substantiate the observed XPS accounts, chemical simulation of the neat electrolyte salts utilizing biphenyl radical anions as a reducing agent were conducted.32 Addition of a single electron (e‒) to a neutral biphenyl molecule generates a radical anion via reduction, with a negative charge and an unpaired e‒ on it (Figure 6). Further injection of an e- to the biphenyl radical anion results in the formation of biphenyl di-radical anion. Biphenyl radical and di-radical anions are estimated to have reduction potentials of ca. 0.40 V and 0.15 V vs. Li/Li+, respectively, and thus, they can be used as powerful

Figure 6. Chemical simulation of the electrochemical stabilities of LiFTFSI, LiFSI and LiTFSI salts.



reducing agents. Figure 6 clearly depicts that while LiFSI and LiFTFSI get easily reduced starting at 1.1‒1.0 V vs. Li/Li+, LiTFSI presented huge resistance and can only get reduced at a very low potential close to that of metallic Li° (< 0.15 V vs. Li/Li+) as proved using biphenyl di-radical anion. Thus, this chemical mock-up corroborates the above XPS results (see Figure 5), but care should be taken considering the different cycling temperature (i.e., 25 and 70 oC for DME- and PEO-based cells, respectively) and physical properties of DME vs. PEO, which may lead to different kinetics in electrochemical reactions and thereby a slightly deviated morphological and chemical information between DME- and PEO-based electrolytes. Based on the XPS results and chemical simulation tests, the SEI layer in the three salts-based electrolytes can schematically illustrated as shown in Figure 7, which clearly demonstrates that while the SEI layer in LiTFSIbased electrolyte is mainly dominated by solvent-derived organic species, the one from LiFSI is richer in inorganic species. LiFTFSI-based electrolyte lies in the middle with an optimum organic-inorganic balance, endowing it with a high mechanical stability, and Li+ cation conductivity. Stability of the salt anions against polysulfide species. As one of the most important requirements for lithium salts used in Li-S batteries, the salt anions should be resistant towards the aggressive polysulfide species. Otherwise, the chemical reaction between them may result in several detrimental impacts on the cycling performance, such as i) deteriorated capacity retention due to the loss of active materials in the S cathode; ii) declining rate capability due to the decrease of ionic conductivity in electrolytes.[4] Figure 8a-d compares the appearance of 1 M LiX/DME (X = FTFSI, FSI and TFSI) solution before (Figure 8a1-d1) and after the addition of PS at room temperature for 60 hours (Figure 8a2-d2). In contrast to the almost unchanged color of LiTFSI/DME after

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adding PS (Figure 8c, d), the greenish colors turn to be light yellowish for LiFTFSI/DME and intense yellowish for LiFSI/DME electrolytes (see Figure 8a and b). This appearance is supported by the diminished ultraviolet visible (UV-Vis) absorbance in the range of 495−570 nm for green light for LiFTFSI- and LiFSI-based solutions (Figure 8e). These results strongly confirm the existence of the chemical reaction between those two anions with PS species. In view of the labile S‒F bond in FTFSI− and FSI−, several intermediates are proposed and computed with density functional theory (DFT) calculations, as shown in Figure 8f. The corresponding Cartesian coordinates of the optimized structures are given in Table S3. The nucleophilic reaction of FTFSI− (or FSI−) with Li2S6 may lead to the substitution of S‒F with S-Sn linkage, where the corresponding reaction energies (∆E, kJ mol−1) reach −359 and −654 kJ mol−1 for FTFSI− and FSI−, respectively. In contrast, the cleavage of S‒C bond within the structure of TFSI− in the presence of PS is significantly less exothermic, with a reaction energy of −221 kJ mol−1. This further confirms that the reactions between anions and PS are more likely to take place for FTFSI− and FSI−. Besides, the substitution of two S‒F bonds in FSI− by PS gives an unstable triple anion (∆E = −654 kJ mol−1) which may also have a low solubility in ether solution as well as in PEO. Then, further decomposition of the trivalent anion irreversibly consumes PS species, thus having a negative impact on the cycling performance of the LiFSI-based cells (see Figure S7 for schematic illustration and Figure 9 for cell performance). In comparison, the CF3SO2N(−)SO2-S6(−) intermediate resulted from the reaction of FTFSI− with PS may be electrochemically reversible and sufficiently soluble in PEO due to its asymetrical structure, thereby leading to an enhanced cycling and rate capability performance in Li-S cells.

Figure 7. Schematic illustration on the SEI layer formed on Li° electrode in the three salts.


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Figure 8. Stability of salt anions against polysulfide species: (a-d) the appearance of 1 M LiX/DME (X = FTFSI, FSI and TFSI) and blank DME solution before and after the addition of PS at room temperature for 60 hours. (e) Normalized ultraviolet-visible (UV-vis) absorption spectra of the PS-added solutions. (f) DFT calculations for the proposed intermediates. Red, yellow, light blue, dark blue, grey, and pink balls in (f) stand for O, S, F, N, C, and Li atoms, respectively. Electrochemical Performance. Though superior physicochemical characteristics and interfacial compatibilities of electrolytes are basic inevitabilities, their practical suitability relies on the proof of concept evaluation of Li-S cells. As manifestation, the rate capability, discharge/areal capacities, cyclability and coulombic efficiency of PEO-based Li-S cells, utilizing LiFSI, LiFTFSI and Li(FSI)0.5(TFSI)0.5 (i.e., 0.5 mole LiFSI + 0.5 mole LiTFSI), have been evaluated at 70oC. The Li-S cell using LiTFSI/PEO electrolyte severely suffers from the shuttling of polysulfide and could not be charged successfully even in the first few cycles (Figure S8). Hence, the equimolar mixture of LiFSI and LiTFSI was used for comparing the difference between molecular level tailoring (i.e., LiFTFSI) and macroscopic level mixing (i.e., Li(FSI)0.5(TFSI)0.5). Prior to the investigation of the anion effect on the electrochemical behavior of Li-S cells, the salt content of LiFTFSI is scrutinized. In contrast to the Li-S cell using LiFTFSI/PEO at EO/Li = 20, the discharge capacity, stability and coulombic efficiency tend to be worse for lower concentrations of LiFTFSI (e.g., EO/Li = 32 and 64, Figure S9). This could be ascribed to the small volume of SEI layer on Liº surface and also lower ionic conductivity (Figure S3). Hence, the polymer to salt ratio has significant influence on the Li-S performance, and thereby the opti-

mal value is fixed as EO/Li = 20 for demonstrating the anion effect in the following discussion. Figure 9a-c presents the discharge profiles of the three electrolytes for the 1st (0.05C), the 15th (0.2C) and the 30th (0.1C) cycles. The corresponding rate capability and areal capacities of the three electrolytes are illustrated in Figure 9d, and the corresponding coulombic efficiency vs. number of cycles is given in Figure S10. Noteworthy, Li-S cell with LiFTFSI-based electrolyte can deliver a very high reversible capacity, ca. 1392 mAh gsulfur–1 in the first cycle (83.2% of the theoretical capacity); yet, LiFSI- and Li(FSI)0.5(TFSI)0.5-based cells can only deliver about 800 and 840 mAh gsulfur–1 in the 1st discharge, which corresponds to 53.3% and 56% that of LiFTFSI-based one respectively. As can be seen from Figure 9d and Figure S3, the LiFTFSI-based electrolyte shows superior capacity at all the tested rates over the 50 cycles with stable coulombic efficiency close to 100%. For instance, it provides a discharge capacity more than two times higher (ca. 510 vs. 240 mAh gsulfur–1) than that of LiFSI and Li(FSI)0.5(TFSI)0.5 based electrolytes at 0.5C, with a good capacity recovery at 0.1C. Besides, an areal capacity of 0.5 mAh cm‒2 can be reached at 0.5C for the LiFTFSI-based cell, whereas only less than a half of it for LiFSI and Li(FSI)0.5(TFSI)0.5-based ones. Effectively, the physical mixture of LiFSI and LiTFSI



causes a poor columbic efficiency and low cycling stability in the corresponding Li-S cell after 35 cycles (Figure S10). Moreover, the superior capacity retentions at a discharge/charge rate of 0.1/0.1C for LiFTFSI- and LiFSIbased cells are comparatively displayed in Figure 9e. LiFTFSI-based cell can deliver ca. 800 mAh gsulfur–1 after 60 cycles, nearly 53% of its initial capacity, while LiFSI-based cell maintains a low discharge capacity of only 450 mAh gsulfur–1 after the 60th cycle. This superior cycling stability can be further extended to higher C-rate at0.5C (Figure 9f), where still a reasonable capacity retention is observed after 100 cycles for the cell using the LiFTFSI/PEO

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electrolyte. The highly impressive discharge capacity associated with high sulfur utilization for LiFTFSI-based LiS cells could be attributed to: i) reversible reaction with PS leading to the formation of highly electroactive intermediate, enhancing the electrochemical kinetics of S cathode; ii) high wettability due to its asymmetrical structure; iii) highly robust anode SEI layer associated with lower resistance and high mechanical stability; (iv) higher wettability of molten-state associated to its flexible and asymmetric structure.

Figure 9. (a−c) Discharge/charge profiles of the Li–S cells using LiX/PEO (X = FTFSI, FSI and (FSI)0.5(TFSI)0.5) electrolytes at 70 oC. (d) Rate capability of those three electrolytes. (e) Cycling stability of the Li–S cells using LiX/PEO (X = FTFSI and FSI) electrolytes. (f) Cycling performance of the LiFTFSI/PEO-based cell at a discharge/charge rate of 0.5/0.5C. (g) Comparison of the cycling performance for ASSLSBs reported in literature and the results of this work. The numbers in squares correspond to the references listed in Table S4 and color code corresponds to type of electrolyte.


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Recently, the use of binary-salts (e.g., LiFSI + LiTFSI) has been reported as a strategy to explore the positive attributes of the individual electrolyte components in the design for high performance liquid Li-S cells.33 However, our work undoubtedly proves that molecular level tailoring outweighs the simple physical mixing as seen from the above results, and the design of salt anions can confer those salts beneficial but antagonistic properties inherited from their predecessors.

Supporting Information. DSC, ionic conductivity, XPS data of SEI components and XPS survey spectra, Cartesian coordinates of the optimized structures, and the cycling performance of all solid-state Li-S cells from literature and Coulombic efficiency of the Li-S cells using the prepared polymer electrolytes (PDF). This material is available free of charge via the Internet at http://pubs.acs.org.

To highlight the impressive results of this work, we compared the performance of LiFTFSI/PEO-based electrolyte in Li-S with literature values (Figure 9g and Table S4). Though without any modification in the Li° anode, S cathode, and no electrolyte optimization (e.g., additives, fillers etc), our results demonstrate to surpass most of the available literature values in terms of areal capacity (i.e., gravimetric energy density).

Corresponding Author

AUTHOR INFORMATION * E-mail: [email protected] (H.Z.). * E-mail: [email protected] (M.A.).

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT

In complement, since elemental sulphur starts to sublime at elevated temperature (e.g., 95 oC for α-S8 (s) cyclo-S8 (g)34), the self-discharge of PEO-based Li-S cells could be more prominent compared to those liquid-based one. Therefore, room temperature conductive SPEs (such as comb-type polymers studied by our group recently35) would be highly justified for the future work in order to lower down the operational temperature of ASSLSBs and mitigate the side reactions raised by elevating the temperature.

This work was supported by the Ministerio de Economía y Competitividad (MINECO) of the Spanish Government through Proyectos I + D Retos program (ENE2015-64907-C21-R and ENE2016-81020-R grants). We also acknowledge funding by the Basque Government through the GVELKARTEK-2016 program. X.J. thanks Basque Government for PhD funding, C.L. thanks the Spanish Government for the Juan de la Cierva scholarship (Ref: FJCI-2015-23898). L.M.R. and H.Z. thank the Basque Government for the Berrikertu program (1-AFW-2017-4 and 1-AFW-2017-2). We are also grateful for computer resources to the i2BASQUE academic network and SGI/IZOSGIker UPV/EHU (Arina cluster).

4. CONCLUSION

REFERENCES

ASSLSBs utilizing LiFTFSI/PEO, as a noble polymer electrolyte, are evaluated and presented. In the first place, the electrolyte showed a highly robust and conductive passivation layer compared to LiFSI and LiTFSI based ones. As a proof of concept, Li-S cells with LiFTFSI/PEO electrolyte presented superior interfacial stability against Li° and of lower resistance, extremely high specific discharge capacity of 1394 mAh gsulfur–1 and high areal capacity of 1.2 mAh cm–2 at the first cycle, high coulombic efficiency (close to 100%), as well as superior rate capability (~ 510 mAh gsulfur–1 at 0.5C). Compared to the electrolyte containing LiFSI or physical mixture of LiFSI and LiTFSI, the high performance of LiFTFSI-based electrolyte is associated to a number of synergistic effects resulted from its molecular structure. LiFTFSI-based electrolytes are proved to boost the properties of the Li° anode and S cathode compartments, an anomalous behavior considering the complexity in the Li-S chemistry.

(1) (2) (3)

Currently, the authors are conducting further tests enlisting with electrolyte optimization (including well thought additives, fillers, etc.), Li° anode and cathode modification etc., so as to further enhance the cyclability and coulombic efficiency of the LiFTFSI-based electrolyte. It is anticipated that this work will shed light on the importance of salt anions and accordingly open a new avenue towards the design of new and tailored SPEs for applications in high performance Li-S cells and beyond.

ASSOCIATED CONTENT

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Figure 1. Comparison of the chemical structures for the LiFTFSI, LiFSI, and LiTFSI salts. 337x250mm (300 x 300 DPI)


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Figure 2. Physicochemical and electrochemical properties of LiX/PEO (X = FTFSI, FSI and TFSI) electrolytes. (a) Opti-cal image of LiFTFSI/PEO membrane. (b) XRD patterns and (c) DSC traces of the three electrolytes. (d) TGA traces for the three electrolytes (top), and their neat components (bottom), (e) Arrhenius plots of ionic conductivity and (f) an-odic stabilities at 70 oC for the three electrolytes. 320x360mm (300 x 300 DPI)



Figure 3. Electrochemical behavior of Li° electrode in the as-prepared electrolytes. Galvanostatic cycling of Li° symmet-ric cells at 0.1 mA cm−2 (half cycle time 2 h) using (a) polymer electrolytes at 70 oC and (b) liquid electrolytes at 25 oC. (c) Galvanostatic polarization of Li symmetric cells at 0.1 mA cm−2 for the liquid electrolytes at 25 oC. 320x260mm (300 x 300 DPI)


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Figure 4. SEM images of Li deposited onto Cu substrates at 0.1 mA cm‒2 (plating time 20 h) in LiX/DME (X = FTFSI (a, d), FSI (b, d) and TFSI (c, f)) electrolytes at 25 oC. 462x287mm (300 x 300 DPI)



Figure 5. (a) XPS spectra of Li° deposited onto Cu substrates at 0.1 mA cm‒2 (plating time 20 h) in LiX/DME (X = FTFSI, FSI and TFSI) electrolytes at 25 oC. (b) Atomic concentration of Li, S, F and N on the Li° deposits as a function of etching time using the three liquid electrolytes. In the figure, R1C‒O and R2C‒O refer to H2C=HC‒O‒CH3 and CH3O‒ respectively. 240x358mm (300 x 300 DPI)


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Figure 6. Chemical simulation of the electrochemical stabilities of LiFTFSI, LiFSI and LiTFSI salts. 230x140mm (300 x 300 DPI)



Figure 7. Schematic illustration on the SEI layer formed on Li° electrode in the three salts. 299x150mm (300 x 300 DPI)


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Figure 8. Stability of salt anions against polysulfide species: (a-d) the appearance of 1 M LiX/DME (X = FTFSI, FSI and TFSI) and blank DME solution before and after the addition of PS at room temperature for 60 hours. (e) Normalized ultraviolet-visible (UV-vis) absorption spectra of the PS-added solutions. (f) DFT calculations for the proposed inter-mediates. Red, yellow, light blue, dark blue, grey, and pink balls in (f) stand for O, S, F, N, C, and Li atoms, respectively. 381x306mm (300 x 300 DPI)



Figure 9. (a−c) Discharge/charge profiles of the Li–S cells using LiX/PEO (X = FTFSI, FSI and (FSI)0.5(TFSI)0.5) electro-lytes at 70 oC. (d) Rate capability of those three electrolytes. (e) Cycling stability of the Li–S cells using LiX/PEO (X = FTFSI and FSI) electrolytes. (f) Cycling performance of the LiFTFSI/PEObased cell at a discharge/charge rate of 0.5/0.5C. (g) Comparison of the cycling performance for ASSLSBs reported in literature and the results of this work. The numbers in squares correspond to the references listed in Table S4 and color code corresponds to type of electrolyte. 274x340mm (300 x 300 DPI)


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Ultrahigh Performance All Solid-State Lithium Sulfur Batteries: Salt

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