Compared to the electrochemical performance of LIBs and NIBs, NaV3O8 nanobelt cathode materials in ZIBs have shown excellent electrochemical performance, including high specific capacity of 421 mA h g−1 at 100 mA g−1 and good cycle stability with a capacity retention of 94% over 500 cycles at 5 A g−1. The good diffusion coefficients and high surface capacity of NaV3O8 nanobelts in ZIBs were in favor of fast Zn2+ intercalation and long-term cycle stability.NaV 3O8 nanobelts electrode for ZIBs shows excellent electrochemical performance, including high specific capacity of 421 mA h g−1 at 100 mA g−1, good rate performance and cycle performance.Compared to the electrochemical performance for LIBs and NIBs, NaV Sodium-ion batteries have been brought into focus for renewable energy and smart grids because of the abundant sodium resources and low cost.0 γ-MnO2  and ZnMnnO2 , high rate capability (120 mA h g−1 at 10 A g−1) and long-term cycling stability (a retention of 94% at 5 A g−1 after 500 cycles).Herein, we develop a facile hydrothermal reaction and subsequent thermal transformation to obtain NaV2.4VO3 and 0.053 g Na2CO3 were added together to a beaker with 50 mL of distilled water and the solution was stirred at 90 °C for 20 min. Then, put the beaker into cool water to room temperature. 1 mL 30% H2O2 was dissolved in the solution and stirred about 2 min with the solution color changed to faint yellow. The solution was adjusted to pH = 1.9 by the titration of 2 mol L−1 HCl and kept for stirring about 30 min, the solution color turns to reddish brown gradually. Finally, the solution was thrown into a 50 mL autoclave and subjected to hydrothermal treatment at 200 °C for 24 h. And then, the autoclave was naturally cooled down and red precipitates were collected, followed with wash with distilled water and alcohol for several times, and finally dried at 60 °C. At last, the product was also annealed in air atmosphere at 400 °C for 4 h with an increasing rate of 3 °C min−1.All chemicals were of analytical grade and were used as received without further purification. In a typical procedure, 0.24 g NaOH, 0.234 g NHλ = 1.5406 Å). The morphology and structure of the samples were characterized by using a scanning electron microscope  and a high resolution transmission electron microscope  operated at 200 kV. X-ray photoelectron spectra (XPS) measurements were conducted using an ESCALAB250 to investigate the elemental composition using an Al Kα source. The nitrogen adsorption–desorption isotherms at the temperature of liquid nitrogen (77 K) were measured on VSorb 4800P Surface Area and Pore Porosimetry Analyzer  with prior degassing under vacuum for 4 h at 150 °C. Total pore volumes were determined using the adsorbed volume at a relative pressure of 0.99. The multi-point Brunauer–Emmett–Teller (BET) surface area was estimated from the relative pressure range from 0.05 to 0.2.The crystallographic structure and phase purity of the as-prepared products were measured by a powder X-ray diffraction analyzer  with Cu Kα radiation  were blended with acetylene black  and polyvinylidene fluoride . The resultant slurry was pasted on an Al foil as cathode electrode and then dried in a vacuum oven. For zinc coin cells, the cathode electrode was fabricated by pressing a mixture of 70% active material, 20% conductive carbon and 10% poly tetra fluoroethylene (PTFE) onto a carbon paper and dried under air at 60 °C for 12 h. A Whatman grade glass fiber was chosen as separator between cathode and anode for Na ion batteries and aqueous zinc ion batteries, respectively. 1 mol L−1 lithium hexafluorophosphate (LiPF6) dissolved in ethylene carbonate (EC), dimethyl carbonate (DMC) and ethyl methyl carbonate (EMC)  was used as electrolyte. For Na coin cells, the electrolyte was 1 mol L−1 NaClO4 dissolved in a solvent mixture of ethylene carbonate (EC) and propylene carbonate (PC) (1 : 1 v/v) and 5% FEC. For zinc ion batteries, the coin cells were assembled under room temperature and 1 mol L−1 ZnSO4·7H2O and 0.2 mol L−1 Na2SO4 mixed solution was used as the electrolyte. Cyclic voltammetry (CV) was carried out using Bio-Logic VSP-300 multichannel electrochemical workstations.Galvanostatic discharge/charge measurements were carried out using Land-2100 automatic battery tester . Electrochemical tests were performed by using a coin battery cell, where a metallic lithium foil, sodium foil and zinc foil were used as the anode electrode for Li/Na ion batteries and aqueous zinc ion batteries, respectively. For Li/Na coin cells, NaV3.3O8 (hereafter denoted as NVO) products are showed in 1.1V3O7.9 phase (JCPDS no. 45-0498), and no other impurities are detected. Nitrogen adsorption–desorption isotherms characterization  scale" fill="currentColor" stroke="none"><path d="M0 440 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z M0 280 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z"/></g></svg>O on the surface, respectively.XRD patterns of the as-synthesized NaVrization  revealed3O8 nanobelts can be described by the following reactions:In our experiment, multi-step synthesis of NaV2(O2)2]3− can be formed by added H2O2 with the solution color turned to faint yellow. For the third step, when the pH value of the solution decreased to 1.9, some [VO2(O2)2]3− can be converted to [V(O2)]3+ with the solution changed to reddish brown. As we known, [VO2(O2)2]3− can be converted to [V(O2)]3+ totally when the pH value is 1.0.The reaction 3O8 and the spot in the SAED pattern was parallel to the length of the NVO nanobelts, indicating the growth direction of the material.The SEM image in −1 with Li in the voltage range from 4 V to 1.5 V. It can be observed that the discharge plateaus located at around 3.2, 2.7 and 2.4 V, implying a complex multiphase transition mechanism during the Li-ion insertion process. The initial discharge capacity of NVO nanobelts is about 200 mA h g−1, which corresponds to a 2.26 electron transfer reaction. A high charge capacity of 195 mA h g−1 can be obtained, indicating a high reversibility of NVO nanobelts cathode materials for Li insertion/extraction reactions. After 50 cycles in −1 with a capacity retention of 76.5% and high coulombic efficiency of 99%. −1, 200 mA g−1, 400 mA g−1 and 800 mA g−1, respectively. The NVO nanobelts electrode presents a high discharge capacity of 76 mA h g−1 at the current density of 800 mA g−1. When the current density finally returned to 100 mA g−1, a capacity of 177 mA g−1 could be recovered, indicating NVO nanobelts remains stable after high rates. In addition, the electrochemical performance of NVO nanobelts as cathode materials in the voltage range from 4 V to 1.5 V for Na-ion batteries was also investigated  and oxidation (Zn2+ extraction) of NVO electrode. The redox couple at about 0.80/1.01 V in the cathodic/anodic sweep represents the redox couple at the equilibrium voltage V5+/V4+ in NVO electrode. Similarly, the redox couple at 0.46/0.73 V represents the V4+/V3+ couple. The multiphase vanadium redox couples commonly induce multistep electrochemical intercalation/extraction of various charge carriers in the vanadium-based layered structure. The second and third discharge profiles are similar and slightly different from the first, indicating a stead structure after first cycle. The largely overlapped CV curves for the NVO electrode suggest the good reversibility of the zinc ions interaction/extraction process and the good retention of crystallinity.Recently, vanadium-based materials have received extensively attention as cathode for aqueous ZIBs with excellent electrochemical performance. Herein, the electrochemical performance of NVO nanobelts was also investigated. The first three cycled electrochemical Zn-ion storage performance of NVO electrode in −1 in the voltage range from 0.2 V to 1.6 V  for NVO. In the second and third discharge process, these two voltage plateaus of NVO remained in every profile. This phenomenon suggested that the electrochemical reactions might occur without irreversible crystal structure change in the subsequent cycles. The rate performance of NVO electrodes is also tested in −1 at current densities of 0.1, 0.5, 1, 2, 5 and 10 A g−1, respectively. When the rate is returned to 1 A g−1, the electrode is able to recover a reversible capacity of 270 mA h g−1, indicating a high reversibility of NVO nanobelts cathode materials for Zn insertion/extraction reactions.Galvanostatic first charge–discharge profiles of NVO electrodes with Zn are examined at a current density of 100 mA gto 1.6 V . In agre−1, respectively. After 50 cycles, the discharge capacities are reduced from 400 to 253 mA h g−1, corresponding to the capacity retention of 62%. −1. The increase specific capacity of NVO from 139 mA h g−1 to the maximum capacity of 163 mA h g−1 for about the initial 20 cycles is attributed to the activation. After reaching a steady state, the electrode exhibits a decrease in capacity, and at the end of 500 cycles, achieves a capacity of 131 mA h g−1 with retention of 94% of the first discharge capacity, indicating the outstanding rate ability and stable cycling ability of NVO nanobelts electrode.−1 at the specific power of 71 W kg−1 based on the active mass of the cathode material, which surpasses those cathodes such as Zn0.25V2O5·nH2O,18 V2O5·nH2O,19 Zn2V2O7,24 Na2V6O16·3H2O,28 Na3V2(PO4)3 3  and Ca0.−1 with Zn were shown in 33i and ν are the peak current (mA) and the corresponding scan rate (V s−1), respectively; a and b are variable parameters in which the latter varies in the range of 0.5–1.0. For a given system, the b value of 0.5 depicts the process of diffusion-limited and the value is 1.0 indicating a capacitive process. Based on the relationship between log(i) and log(ν) for two different peaks, the b-values determined by the slopes of the four redox peaks are 0.61 and 0.79  are calculated as given in −1, 61% of the capacitive contribution is capacity-limited. When the scan rate improved to 1 mV s−1, nearly 83% of the capacitive contribution is capacity-limited. The results indicate that the capacitance effect play an increasingly important role as the scanning rate is increased. The electrochemical kinetics of NVO indicates that the surface dominated capacitive behaviors of extrinsic NVO nanobelts are in favor of fast Zn2+ intercalation/extraction.From 2+ ions in NVO. The GITT data were collected at a current density of 100 mA g−1 for 10 min and a rest interval of 30 min  measurements are carried out to better understand the kinetics during the diffusion of Znf 30 min . The ion cm2 s−1 . The difex situ XRD analysis in −1. The obvious reflection peak at 2θ = 12.7° (d = 0.69 nm) as attributed to the (002) plane of NVO. When the electrode was discharged to 0.8 V, a few new peaks appeared at 8.0°, 16.1°, 24.5°, 26.6°, 28.7° and 32.9°, which could be indexed to a triclinic Zn4SO4(OH)6·5H2O phase (JCPDS no. 39-0688). The (002) plane of NVO becomes weaker and slightly shifted toward higher 2θ positions from 12.7° to 12.9°. With more Zn2+ intercalated into the electrode from 0.8 V to 0.2 V, the reflection peaks of Zn4SO4(OH)6·5H2O are gradually strengthened, and the (002) plane of NVO continuously shifted to a higher 2θ position (13.0°), which corresponded to an enlarged interlayer distance of 0.67 nm. This reduction in the interplanar spacing from 0.69 to 0.67 nm indicates an improvement in the structural coordination owing to the strong electrostatic interaction between the intercalated Zn2+ and the (V3O8)− layers.28 During the subsequent charging process, the reflection peaks of Zn4SO4(OH)6·5H2O gradually disappeared. The formation of this phase can be described by the following reactions:37Furthermore, θ position from 13.0° to 12.5°, which demonstrated that the Zn2+ intercalation and extraction in NVO was highly reversible. The XRD of the electrode after 500 cycles at a current density of 5 A g−1 is investigated in 3O8 electrode.Meanwhile, the (002) plane of NVO shifted back to a lower 24.3O8 nanobelts as a promising cathode material for LIBs, NIBs and ZIBs, respectively. Compared to the electrochemical performance for LIBs and NIBs, NaV3O8 nanobelts cathode materials for ZIBs have shown excellent electrochemical performance, including high specific capacity of 421 mA h g−1 at 100 mA g−1 and good cycle stability with a capacity retention of 94% over 500 cycles at 5 A g−1. Ex situ XRD showed that the Zn2+ intercalation and extraction in NaV3O8 nanobelts was highly reversible. The capacitive behavior and good Zn ions diffusion coefficients played an important role in fast Zn2+ intercalation/extraction and long cycle stability. This work also indicates that NaV3O8 nanobelts have a great potential as a cathode material for ZIBs in large-scale electric energy storage applications.In this work, a facile hydrothermal reaction was developed to prepare NaVThere are no conflicts to declare.


The previouslyreported monoclinic form of the 1:1 cocrystal of LIS and PRO, LISPRO(α),and a new thermodynamically stable orthorhombic polymorph, LISPRO(β),were prepared and characterized. The two polymorphs form square grid,sql, topology coordination networks and differ mainly in the conformationof the salicylate ions and positioning of the sql nets. LISPRO(α)was observed to transform to LISPRO(β) under slurry conditions.The 1:1 ionic cocrystal of L4M and PRO (L4MPRO) was found to formthree polymorphs. Apart from the previously reported orthorhombiccrystal form, L4MPRO(α), two new monoclinic crystal forms, L4MPRO(β)and L4MPRO(γ), were obtained by modifying crystallization conditions.The new polymorphs were found to be metastable, undergoing transformationsto L4MPRO(α) upon exposure to humidity. Experimental conditionsthat induce transformations between the polymorphs of LISPRO and L4MPROare detailed, and the structural differences between the polymorphsare discussed in the broader context of polymorphism.The occurrence ofpolymorphism in ionic cocrystals formed by twolithium salts, lithium salicylate (LIS) and lithium 4-methoxybenzoate(L4M), and  l-proline (PRO)as a coformer was investigated, affording a more thermodynamicallystable polymorph of lithium salicylate l-proline [LISPRO(β)]and two metastable polymorphs of lithium 4-methoxybenzoate l-proline (L4MPRO). The crystal structures and experimental conditionsthat enabled transformations between polymorphs are detailed.Polymorphism in ionic cocrystals of lithiumsalicylate (LIS)and lithium 4-methoxybenzoate (L4M) with  This therapeutic window hinders the use of lithiumas a treatment as it requires regular and intensive evaluation ofblood, thyroid, and kidney function to monitor the effects of lithiumtoxicity. Lithium toxicity can cause loss of consciousness, epilepticseizures, constipation, and hyperreflexia, which can negatively impacta patient’s compliance with treatment.7 Despite these drawbacks, lithium remains the most viable optionfor several indications, so formulation strategies to enable controlleddelivery, reduce dose frequency, and decrease peak serum lithium concentrationshave been studied to minimize side effects. In this context, our group6 reported the lithium salicylate (LIS) l-proline ionic cocrystal, LISPRO, which is currently being developedunder the trade name LiProSal. Importantly, LISPRO was found to exhibitlower peak serum lithium concentration and longer retention time inthe blood stream than lithium carbonate. In principle, LISPRO shouldtherefore reduce lithium toxicity and enable a less-demanding dosageschedule. Additional advantages and mechanisms of action for LISPROhave been elaborated by Habib et al.11 The related ionic cocrystal (ICCs),lithium 4-methoxybenzoate l-proline (L4MPRO), disclosed ina patent,12 offers similar promise as atherapeutic. As noted by a reviewer of this article, there is someambiguity as to whether ICCs of the type studied herein should bereferred to as coordination polymers. In our opinion, these termsare not mutually exclusive and the context would affect which termis adopted. From the perspective of pharmaceutical science, the termionic cocrystal would therefore be appropriate.14Lithium salts have a long history of therapeutic use in the treatmentof bipolar disorder,15 this group of compounds remainsunderstudiedin the literature from a crystal engineering perspective.16 Forty-seven examples20 of natural amino acid-based lithiumICCs are archived in the Cambridge Structural Database21 , and only one case of lithium ICC polymorphismhas been reported thus far.5 Polymorphism,26 the phenomenon in which a chemical compound exhibits different crystalstructures, is long known but understudied outside pharmaceuticalcompounds. In this context, the United States Food and Drug Administration(FDA) has recognized the relevance of polymorphism to drug productsand defined polymorphism as “the ability of a substance toexist in at least two crystalline forms with different crystal packingarrangements and/or conformations of molecules in the crystal lattice”.27Despite the evident pharmacological benefits of lithium-basedICCs,28 However, in 1809, Humphrey Davy noted that both diamond and graphitewere made of carbon, with the only difference being the arrangementof carbon atoms in the solid phase.29 Thefirst report of a polymorphic organic compound dates back to 1832with the discovery of a polymorph of benzamide by Liebig and Wöhler,which was obtained during the heat–cool crystallization ofan aqueous solution of benzamide.30The term “polymorphism”was first used in 1822 byMitscherlich when he noted different physical and chemical propertiesof crystals of arsenates and phosphates.32 Further, if the polymorphic landscape of a compound is not understood,the unexpected emergence of an unwanted solid form, or conversionbetween polymorphs, can result in negative consequences, as exemplifiedby Ritonavir.33The primary reason that the pharmaceutical industry is interestedin studying polymorphs is that most active pharmaceutical ingredients(APIs) are delivered as crystalline drug substances and polymorphscan vary in their physical, chemical, mechanical, and biopharmaceuticalproperties as well as their stability, thereby impacting utility.36 screening fornew solid forms of drug substances and studying their stability hasbecome a routine and required aspect at the pre-clinical stage ofdrug development.39 Such screening can include the isolation and studyof multicomponent crystals such as solvates, hydrates, and pharmaceuticalcocrystals42 .43 Pharmaceutical cocrystals,which can offer very different physicochemical properties and a diverserange of crystal forms without compromising therapeutic benefits,46 have been bolstered by the FDA and EMA releasing regulatory guidelinesfor industry on the use of pharmaceutical cocrystals.48 Indeed, the number of marketed drug products based on APIs thatare pharmaceutical cocrystals has increased in recent years,49 with four new drug products approved between2014 and 2017.49 In addition, a new pharmaceuticalICC is the API in Seglentis, which received FDA approval in October2021.51Whereas differencesin polymorphic properties such as aqueous solubilityand melting point are typically negligible,54 has been reported, it remains understudied in thecontext of pharmaceuticalcocrystals. Herein, we report the synthesis and characterization ofpolymorphs of ICCs of LIS and L4M with l-proline. Only oneform of each ICC had been reported prior to this study, LISPRO(α)6 and L4MPRO(α).55Whereas polymorphism in cocrystals22.1l-proline, lithium hydroxide, and 4-methoxybenzoicacid) or Alfa Aesar (LIS) and used without further purification. PowderX-ray diffraction (PXRD) data were measured on an Empyrean diffractometer equipped with a PIXcel3D detectorusing a Cu Kα radiation source in reflection geometry. Thermogravimetricanalyses were performed on a TA Instrument Q50 TG. Differential scanningcalorimetry (DSC) analyses were carried out on a TA Instrument DSCQ20. Fourier transform infrared (FTIR) spectra were collected on aPerkinElmer Spectrum 100 spectrometer with a Universal ATR accessory.Reagents were purchasedfrom Sigma-Aldrich (2.2α (λ = 0.71073Å) or Cu Kα (λ = 1.5418 Å) radiationon a Bruker D8 Quest fixed-chi diffractometer equipped with a BrukerAPEX-II CCD detector and a nitrogen-flow Oxford Cryosystem attachment.Data was indexed, integrated, and scaled in APEX3.56 Absorption corrections were performed by the multi-scanor numerical method using SADABS57 or TWINABS.58 Space groups were determined using XPREP59 as implemented in APEX3. The SHELX-2014 programpackage, implemented in OLEX260 v1.2.8or v1.3.0, was used for structure solution and refinement. Structureswere solved using the intrinsic phasing method (SHELXT)61 and refined with SHELXL62 using the least-squares method. All non-hydrogen atoms were refinedanisotropically. Hydrogen atoms were placed in calculated positionsfrom the molecular geometry and assigned isotropic thermal parametersthat depended on the equivalent displacement parameters of their carriers.Crystal data were deposited with the Cambridge Crystallographic DataCentre (CCDC 2116205–2116210 and 2142069–2142072). Selected crystallographic data and refinementparameters for the crystal structures refined from data collectedat 100 K are given in Supporting Information.Crystals of LISPRO (α and β) andL4MPRO  were examined under a microscope,and suitable single crystals were selected for single-crystal X-raydiffraction (SCXRD) analysis. Single-crystal structures were determinedat RT or 100 K using either Mo K2.32.3.1l-proline  were dissolved in 2mL H2O and heated to 70 °C for ca. 1 h in an openvial to allow solvent evaporation until crystals had formed. The initiallyformed crystals were harvested and stored under oil. These crystalswere found to be of LISPRO(α), as verified by SCXRD; however,a PXRD of the crystals was found to be a mixture of LISPRO(α)and LISPRO(β)  and ISPRO(β) 1.2.3.2l-proline  in1.5 mL H2O was conducted under ambient conditions for 24h. The resulting powder was filtered and air-dried. Recrystallizationexperiments using various solvents such as MeOH, EtOH, MeCN, and i-PrOHwere conducted in order to obtain single crystals suitable for SCXRD.Then, 14 mg of the sample obtained from slurry was dissolved in 1.5mL EtOH in a loosely covered vial and left to undergo slow evaporation.Colorless needle-shaped crystals of LISPRO(β) were obtainedafter 3 days.A 1:2 water slurryof LIS  and 2.3.3l-proline  in 0.5 mL H2O was conducted at ambientconditions for 24 h, affording colorless needle-shaped crystals ofL4MPRO(α), as verified by a PXRD.L4M was preparedby grinding lithium hydroxide  and 4-methoxybenzoicacid  in a mortar and pestle with 0.44 mL distilledwater until a free-flowing powder was obtained (approximately 15 min)and dried in an oven at 75 °C for 20 h. A 1:10 water slurry ofL4M  and l-proline  with 5 mol % L4MPRO(α) and 0.04 mL distilled waterin a mortar and pestle until a free-flowing powder was formed (approximately15 min). A PXRD performed on the resultant powder revealed it to bea mixture of L4MPRO(α) and L4MPRO(β). Then, 30 mg of thismixture (0.11 mmol) was slurried in 0.15 mL EtOH under ambient conditionsfor 3 days, resulting in L4MPRO(α), as confirmed by PXRD.Seeding-assisted mechanochemicalsynthesis of L4MPRO(α) was performed by grinding L4M  and 2.3.4l-proline  in 2 mL H2O andheating to 70 °C for ca. 1 h in an open vial to allow solventevaporation until crystals had formed. The initially formed crystalswere harvested and stored under oil. These crystals were determinedto be L4MPRO(β) via SCXRD. A powder sample was prepared withL4M  ground in a mortar and pestlewith l-proline  and 0.11 mL distilledwater until a free-flowing powder was formed (approximately 15 min).Alternatively, L4MPRO(β) was prepared by slurrying L4M  and l-proline  for24 h in 0.52 mL EtOH, ethyl acetate, or acetone or 0.17 mL dimethylformamide(DMF), chloroform, or tetrahydrofuran. The slurry procedure with 0.17mL MeCN yielded a mixture of L4MPRO(β) and L4MPRO(γ).Colorless platecrystals were prepared by dissolving L4M and 2.3.5l-proline  weredissolved in a heated mixture of 4 mL MeOH and 0.1 mL H2O. The vial was loosely covered, and after 24 h, large plate-shapedcrystals of L4MPRO(γ) formed along the vial walls. A powderedsample of L4MPRO(γ) was formed by grinding L4M  and l-proline  with0.32 mL dimethyl sulfoxide (DMSO) until a homogenous powder was formed.Additionally, L4MPRO(γ) was obtained by slurrying L4M  and l-proline  ineither 0.52 mL MeOH or i-PrOH for 24 h. A similar slurry procedurewith 0.17 mL MeCN yielded a mixture of L4MPRO(β) and L4MPRO(γ).L4M  and 2.463 The synthesizedICCs were subjected to such stability testing by placing 30 mg ofeach compound in a glass vial loosely covered with aluminum foilin a humidity chamber at 75% RH and 40 °C. Aliquots were removedafter 14 days, and PXRD data were collected.In order to explorethe effect of humidity on lithium ICCs, accelerated stability testingunder conditions used in the pharmaceutical industry (40 °C and75% RH) were employed.2.5Figures S1, S2, and S7–S10.The purity ofthe ICCs was assessed by PXRD and DSC. Comparisons of the experimentalPXRD patterns of as-prepared ICCs with PXRD patterns calculated basedon crystal structures are presented in 33.16 through slurrying afforded PXRD patterns  that did not match the calculated PXRD pattern of the previouslyreported form of LISPRO, herein called LISPRO(α). Interestingly,the PXRD pattern did resemble the experimental PXRD reported by Smithet al.6 Single-crystal analysis of an ethanol-recrystallizedsample of LISPRO revealed a new polymorph, LISPRO(β). The experimentalPXRD of the as-prepared LISPRO and the calculated PXRD patterns ofthe α and β polymorphs are overlaid in S1, and S2. The experimental PXRD pattern of LISPRO(β) isin good agreement with that calculated from the crystal structureof LISPRO(β). Subsequent attempts to obtain a pure sample ofLISPRO(α), either by rapid solvent evaporation or by slurrying,failed, with all samples containing at least some LISPRO(β)according to PXRD data (S1). A close inspectionof the experimental PXRD of the first report of LISPRO6 hints at this same outcome, with a mixture of peaks consistentwith LISPRO(α) and LISPRO(β). This precluded us from conductinga 50:50 slurry of the two LISPRO polymorphs. Rather, a mixture ofLISPRO(α) and LISPRO(β) obtained from heating–coolingwas slurried in EtOH for 24 h. This experiment afforded pure LISPRO(β),indicating that it is more stable than LISPRO(α)  under these conditions. No polymorphictransformation was observed when crystals of LISPRO(β) wereexposed to accelerated stability testing conditions for 14 days . Long-term stabilitystudies of LISPRO(β) under ambient conditions confirmed solid-formstability for at least 6 months.Severalattempts to synthesizethe previously reported form of LISPROXRD data 1 and S1.3.2Figures S7–S10, S16).The previously reported form, L4MPRO(α),55 was obtained as single crystals from the H2Oslurry of l-proline and L4M in a 10:1 molar ratio. Attemptsto prepare L4MPRO(α) through SDG64  failed; a 1:1 water grind yielded L4MPRO(β), while a 1:1DMSO grind afforded a third polymorph, L4MPRO(γ). The 1:1 slurryexperiments also failed to produce L4MPRO(α), with most solvents,for example, ethanol and DMF, resulting in L4MPRO(β), whileMeOH and i-PrOH afforded L4MPRO(γ) . Similar experiments conducted in water generally failedto produce L4MPRO, resulting in L4M.As illustrated in Figures S31–S33). Seeding the as-formed L4MPRO(α) crystals into a 1:1 water-basedSDG synthesis of LM4PRO resulted in a powder comprising L4MPRO(α)with only minor quantities of L4MPRO(β) . These results suggest that whereas LM4PRO(α)is thermodynamically favored, kinetics drives the formation of theother forms. The 50:50 slurries of either the α and βpolymorphs or the α and γ polymorphs were then performedin ethanol, each resulting in transformation to the α-form . Long-term stability studiesof the polymorphs under ambient conditions revealed that L4MPRO(α)and (β) remained stable for at least 6 months while L4MPRO(γ)converted to L4MPRO(α) .That L4MPRO(α) did not formunder slurry conditions wouldtypically indicate relative instability compared to the other polymorphs;however, when L4MPRO(β) and (γ) were exposed to acceleratedstability-testing conditions  for 14 days, transformationto L4MPRO(α) occurred  > L4MPRO(β)> L4MPRO(γ) , while at 100 K, the order is as follows: L4MPRO(α)> L4MPRO(γ) > L4MPRO(β) .The higher thermodynamic stability of L4MPRO(α) isalso consistentwith density.69In the case of thisstudy, slurry-based and mechanochemical approacheswere found to be complementary. Slurrying was informative with respectto relative stability and, from a screening perspective, affordedall polymorphs except for LISPRO(α), while mechanochemistryinitially failed to afford L4MPRO(α). These observations furtherhighlight the benefit of using multiple screening approaches for cocrystaldiscovery.3.3Inthis section, we refer to crystal structures determined from datacollected at 100 K unless stated otherwise.3.3.1P212121. The unit cellwas found to contain eight lithium cations, eight salicylate anions,and eight l-proline molecules, with two formula units [Li(Pro)]in the asymmetric unit (Z′ = 2). Within theasymmetric unit, one salicylate and one l-proline moleculeis disordered, with the aromatic ring of the salicylate moleculesrotated to allow intramolecular hydrogen bonding with either of itscarboxylate group oxygen atoms. The major components of the salicylateand l-proline disorder have site occupancies of 0.833(8)and 0.802(11), respectively. Lithium cations exhibit a tetrahedralcoordination geometry, being linked by four bridging carboxylate moieties,two from salicylate and two from l-proline. The resultingstructure is a zigzag square grid network propagating in the (001)plane . The aminegroups of l-proline zwitterions form hydrogen bonds withinthe square grids to the carboxylate oxygen atoms of l-prolinewith a distance of 2.733 (4) Å or 2.779(4) Å and to thecarboxylate oxygen of a neighboring salicylate with a distance of2.830(4) Å or 3.081(8) Å (for the major component). Additionalhydrogen bonds are formed by salicylate ions between the hydroxylgroup and the carboxylate group, with distances of 2.593(8) and 2.583(4)Å. Salicylate and l-proline molecules form square gridsthat alternate in their orientation. The l-proline zwitterionsalternately point above and below the walls of the zigzag grid, whilesalicylate ions alternate between pointing near-parallel or at anangle to the walls, enabling the close packing of adjacent squaregrids.LISPRO(β) crystallizedin the orthorhombic space group 3.3.26 LISPRO(α) crystallized in P21 with a unit cell containing four lithium cations, four salicylateanions, and four l-proline zwitterions. There are two formulaunits [Li(Pro)] in the asymmetric unit (Z′= 2). One salicylate and one l-proline molecule of the asymmetricunit show a disorder similar to that observed for LISPRO(β).The major components of the salicylate and l-proline disorderhave site occupancies of 0.761(15) and 0.55(2), respectively. Lithiumcations exhibit a tetrahedral coordination geometry through bondingto salicylate anions and l-proline molecules via carboxylategroups, resulting in a square grid propagating in the (001) plane. Amine groups of l-proline zwitterionsform hydrogen bonds within the square grids to the carboxylate oxygenof l-proline with a distance of 2.752(7) or 2.773(7) Åand to the carboxylate oxygen of a neighboring salicylate with a distanceof 2.843(8) or 3.125(15) Å. Additional hydrogen bonds form betweenthe hydroxyl and carboxylate groups of salicylate molecules, withdistances of 2.610(9) and 2.557(15) Å (for the major component).Salicylate and l-proline molecules in the square grids alternatein their arrangement relative to the grid, with l-prolinezwitterions alternating between pointing above and below the wallsof the grid and salicylate ions positioned either in a parallel manneror at an angle, resulting in a similar structure to LISPRO(β).To allow appropriate comparison of structures, SCXRD data for LISPRO(α)was recollected at 100 K. As previously reported,3.3.3P21 and P212121 space groups, respectively). The volume of the unit cell is doubledin the case of the β polymorph due to the different space groupand doubling of the unit-cell parameter c. In bothLISPRO structures, one l-proline molecule is disordered betweentwo positions on the C4 methylene carbon, pointing either above orbelow the plane of the 5-membered ring. In addition, the structureof each polymorph contains one symmetrically independent salicylateion that shows the disorder of its aromatic ring, with major and minorcomponents. In the structure of LISPRO(α), the hydroxyl groupsof disordered salicylate ions (the major component) and a second symmetricallyindependent salicylate orient in opposite directions within the same2D net, that is, in the [1̅00] and [100] directions, respectively.Meanwhile, in LISPRO(β), the hydroxyl groups of both symmetricallyindependent salicylates are positioned in the same direction [1̅00]within the nets (considering the major component of the disorder),with adjacent 2D nets having hydroxyl groups oriented in the oppositedirection [100] .Thesimilarity of PXRD patterns of the LISPRO polymorphs can be explainedby the closely related crystal structures of LISPRO α and β,despite differences in the crystal symmetry (on [100] 4. Conforon [100] 3–5. Alon–3 compared to 1.385 g cm–3 for LISPRO(β). At RT , the relative density ispreserved, with β being less dense (1.334 g cm–3) than α (1.350 g cm–3). However, it shouldbe noted that the quality of the X-ray diffraction data at room temperaturewas affected by the poor diffracting properties of LISPRO crystals,reducing the accuracy of the unit-cell parameters and the calculatedcrystal density.The structural similaritiesbetween LISPRO(α) and LISPRO(β)mean that there is a small difference in the calculated densities.At 100 K, the α polymorph is the denser phase with a calculateddensity of 1.387 g cm3.3.4P212121 with Z′ = 1. SCXRD revealed that L4MPRO(α) contains four lithiumcations, four 4-methoxybenzoate anions, and four l-prolinemolecules in the unit cell, with the asymmetric unit comprising onemolecule of l-proline, one 4-methoxybenzoate anion, and onelithium cation plane . The lithium cations are bridged in the[100] direction by 4-methoxybenzoate ions and in the [011] and [011̅]directions by l-proline zwitterions, in an alternating manner,resulting in a wave-like square grid net. The protonated amine groupsof l-proline zwitterions are H-bonded to a carboxylate oxygenof a neighboring l-proline with a distance of 2.843(2) Åand to the carboxylate oxygen of a neighboring 4-methoxybenzoate anionwith a distance of 2.880(2) Å. The folding of the square gridcan be assessed by the Li–C(carbonyl)–Li angle, as measuredbetween two lithium cations linked by an l-proline zwitterion,and the Li–Li–Li angle between consecutive lithium cationsarranged in a zigzag manner in the [010] direction. For the squaregrid of L4MPRO(α), these angles are 149.5(1) and 103.9(1)°,respectively. The 4-methoxybenzoate moieties lie nearly parallel tothe square grid walls at an angle of 6.5(7)°, as measured betweenthe plane of parallel C6–C7 bonds of 4-methoxybenzoate ionsand the plane formed by lithium ions in one square of the square grid. As all 4-methoxbenzoategroups orient in the same direction within individual square grids,a gap is created between each “rung” of the wave-likenet and is filled by a 4-methoxybenzoate of an adjacent grid .L4MPRO(α) crystallized in the orthorhombic space group m cation 6. Each l3.3.5P21 with Z′ = 1. SCXRD revealed thatL4MPRO(β) contains two lithium cations, two 4-methoxybenzoateanions, and two l-proline molecules in the unit cell, withthe asymmetric unit comprising one molecule of l-proline,one 4-methoxybenzoate anion, and one lithium cation. Like L4MPRO(α),each lithium cation is tetrahedrally coordinated and bridged by fourcarboxylate moieties, two from l-proline and two from 4-methoxybenzoateanions, to form square grid networks in a staircase-like arrangementalong the (001) plane . Within the square grid,protonated amine groups of l-proline zwitterions form H-bondsto a carboxylate oxygen of a neighboring l-proline with adistance of 2.790(2) Å and to a carboxylate oxygen of a neighboring4-methoxybenzoate with a distance of 2.761(2) Å. The Li–C(carbonyl)–Liangle between two lithium cations linked by an l-prolinezwitterion and the Li–Li–Li angle between consecutivelithium cations arranged in a zigzag manner in the [010] directionare 134.2(1) and 105.8(1)°, respectively. Unlike L4MPRO(α),the 4-methoxybenzoate groups are not aligned with the walls of thesquare grid, with an angle of 43.3(9)° between the plane of parallelC6–C7 bonds of 4-methoxybenzoate ions and the plane formedby four lithium ions . Similar to L4MPRO(α), all 4-methoxybenzoate anions pointin the same direction within a square grid, with the resulting gapbeing filled by 4-methoxybenzoate anions of an adjacent 2D net.L4MPRO(β)crystallized in the monoclinic space group 3.3.6P21, however, with Z′= 2. SCXRD revealed that L4MPRO(γ) contains four lithium cations,four 4-methoxybenzoate anions, and four l-proline moleculesin the unit cell with an asymmetric unit comprising two formula units.As with the α and β polymorphs, each lithium cation istetrahedrally coordinated and bridged by four carboxylate moietiesof two l-proline molecules and two 4-methoxybenzoate anionsto form a square grid network in a staircase-like arrangement alongthe (001) plane . It should be noted thatdue to the presence of two formula units in the asymmetric unit, onlylithium cations not related by symmetry are directly bonded to eachother via ligands. Moreover, neither the two l-proline moleculesnor the two 4-methoxybenzoate ions that coordinate with lithium cationsare symmetrically related. l-proline zwitterions form hydrogenbonds within the square grids to the carboxylate oxygen of l-proline with a distance of 2.784(5) or 2.771(5) Å and to thecarboxylate oxygen of a neighboring 4-methoxybenzoate with a distanceof 2.777(4) or 2.785(5) Å. The Li–C(carbonyl)–Liangle between two lithium cations linked by l-proline zwitterionsis 155.6(2) or 155.5(2)°, and the Li–Li–Li anglebetween lithium cations in the square grids is 110.8(2)°. Unlikethe α and β polymorphs, the 4-methoxybenzoate groups inthe L4MPRO(γ) square grids alternate in their position withrespect to the grid walls, between near-parallel orientation [withan angle of 4.5(2)° measured between the plane of parallel C6A–C7Abonds and the plane formed by the lithium cations in the grid wall)]and out-of-plane with an angle of 62.7(2)° . As a result, nocavities exist in the sides of the 2D framework, a situation similarto the LISPRO structures.Similar to L4MPRO(β),L4MPRO(γ) crystallized in the monoclinic space group 4l-proline with LIS and L4M was studiedby applying various experimental protocols that resulted in new polymorphicforms of the ICCs LISPRO and L4MPRO. Two crystal structures, LISPRO(α)and L4MPRO(α), were known prior to this study, and three newcrystal forms were obtained and characterized: the β polymorphof LISPRO and the β and γ forms of L4MPRO. LISPRO(β)and L4MPRO(α) were found to be thermodynamically stable underthe conditions studied. L4MPRO(α) exhibited a higher crystaldensity compared to polymorphs β and γ at both 100 K andRT. Pure samples of LISPRO(β) were isolated by water slurry,while all polymorphs of L4MPRO were isolated from solvent drop grindingand/or slurry. All five crystal structures analyzed herein exhibitsquare grid, sql, network structures. The polymorphs of LISPRO areclosely related in terms of structure and density, with only minordifferences with respect to the conformation of l-prolineand salicylate moieties and relative positioning of the square grids.The differences between the three crystal forms of L4MPRO are morepronounced, with α and β being similar to each other interms of the orientation of the l-proline and 4-methoxybenzoatemoieties and L4MPRO(γ) being distinctly different. To conclude,whereas we herein report polymorphism in two previously reported ICCsusing conventional methods, polymorphism in ICCs remains largely understudied.It is therefore premature to draw broad conclusions about the propensityof ICCs to exhibit polymorphism, and this is a subject that deservesattention through systematic studies.Cocrystallization of


The specific capacity of the optimal electrode was very high (517 mA h gelectrode−1). Furthermore, considering the weights of the anode, separator, cathode, and electrolyte, the Li/S cell exhibited a high practical energy density of 250 W h kg−1. The areal capacity was also high (8.5 mA h cm−2) owing to the high SPAN loading of 16.37 mg cm−2. After the introduction of 10 wt% multi-walled carbon nanotubes as a conducting agent, the SPAN disk electrode exhibited excellent cyclability while maintaining a high energy density. This strategy offers a potential candidate for Li/S batteries with high practical energy densities.Sulfurized polyacrylonitrile (SPAN) is a promising active material for Li/S batteries owing to its high sulfur utilization and long-term cyclability. However, because SPAN electrodes are synthesized using powder, they require large amounts of electrolyte, conducting agents, and binder, which reduces the practical energy density. Herein, to improve the practical energy density, we fabricated bulk-type SPAN disk cathodes from pressed sulfur and polyacrylonitrile powders using a simple heating process. The SPAN disks could be used directly as cathode materials because their π–π structures provide molecular-level electrical connectivity. In addition, the electrodes had interconnected pores, which improved the mobility of Li A simple synthesis procedure to prepare bulk-type SPAN electrodes toward the realization of Li/S batteries with enhanced practical energy densities. However, binders have low electrical conductivities, which inhibits electron transfer and reduces the energy density. Therefore, a binder-free, highly interconnected electrode structure with a minimum amount of inactive materials is required. (ii) Introduce synergism between the electrode and electrolyte. Insufficient sulfur loading and the use of excess electrolyte have failed to meet the demands of practical applications. Therefore, to achieve high practical energy densities, sulfur electrodes should be reasonably designed to maintain sufficient electrochemically active sulfur species at a reasonably high sulfur loading with a minimal amount of electrolyte.Current Li/S batteries exhibit low practical energy densities because a low proportion of sulfur in the electrode and large amounts of electrolyte are used to compensate for the insulating properties of sulfur and the solubility of lithium polysulfides, respectively. To achieve high practical energy densities, the following strategies can be adopted: (i) incorporate a minimum amount of inactive materials, such as binders and conducting agents, into the electrode. In particular, conventional sulfur electrodes are generally based on particle-type materials and prepared 8) as an active material is not suitable because large amounts of ether-based electrolytes and conducting agents are required. In previous reports, more than ∼10 μL of electrolyte has typically been used for 1 mg of sulfur.22–31 The Chen group reported the use of a minimal amount of electrolyte (∼3 μL of electrolyte for 1 mg of sulfur), but a low specific capacity of 590 mA h gsulfur−1 was observed, indicating that this amount of electrolyte still reduces the utilization of sulfur.32As mentioned above, the use of elemental sulfur  in a carbon matrix such as sulfurized polyacrylonitrile (SPAN)36–43 or sulfur-infiltrated specific nanoporous carbon.44,45 Unlike S8, S2–S4 can directly form Li2S; therefore, when using short-chain sulfur as an active material, the formation of long-chain lithium polysulfides can be ignored.46 Consequently, smaller amounts of carbonate-based electrolytes can be used, resulting in enhanced electrochemical performance. Thus, electrically connected bulk SPAN should be synthesized to improve the practical energy densities of Li/S batteries.Another approach is to adopt carbonate-based electrolyte, in which lithium polysulfides are insoluble.+ ion mobility. In addition, the electrolyte/sulfur ratio was reduced by using a carbonate-based electrolyte with high sulfur utilization . To determine the optimal preparation conditions for the SPAN disk electrodes, different heating conditions and synthesis methods were investigated. The obtained bulk-type SPAN electrodes exhibited high practical energy densities and good electrochemical performance. In addition, the reaction mechanism of the bulk-type SPAN disk was clarified and the electrochemical properties were improved by adding a minimal amount of conducting agents.Herein, we designed a binder-free and interconnected electrode using a simple compacting and heating process. The obtained bulk-type free-standing sulfur disk electrodes were composed only of SPAN, unlike conventional particle-type electrodes. Therefore, the electrodes had interconnected pores, which improved the Li2.2.1.Mw = 150 000, Sigma-Aldrich) powders were mixed at a weight ratio of 80 : 20 using a mortar and pestle. The mixed powder was compressed by applying a pressure 500 kg in a cylindrical mold with a diameter of 10 mm to form pellets (denoted SP000). The SP000 pellets were covered with Al foil to prevent the evaporation of sulfur at high temperatures. Then, the samples were heat-treated at 150, 300, 450, or 600 °C for 6 h in an argon atmosphere to obtain free-standing bulk-type electrodes . To determine the effect of compression, SP450 was ground using a mortar and pestle. Then, the SP450 powder was compressed under the same conditions, and the resulting sample was denoted SP450compact.Sulfur  and polyacrylonitrile  and β-cyclodextrin (Sigma-Aldrich) at a weight ratio of 80 : 10 : 10 using distilled water as a dispersant by ball milling at 300 rpm for 3 h. The resulting slurry was coated on Al foil and then dried at 60 °C overnight, and the obtained sample was denoted SPconventional. To improve the performance of SP450, MWCNTs were added to SP000 (sulfur/PAN/MWCNT ratio of 70 : 20 : 10). This mixture was treated under the same conditions used to prepare SP450, and the obtained sample was denoted SP450-C.2.2.The crystalline structures of the electrodes were confirmed by X-ray diffraction . The thermal decomposition behaviors of the raw materials and the SPpellet samples were determined by thermogravimetric analysis . The surface morphologies were examined using field emission scanning electron microscopy  coupled with energy-dispersive X-ray spectroscopy . The sulfur contents in the electrodes were determined using elemental analysis . The resistance values of the SPpellet samples were measured using a multimeter inside a glovebox. The molecular structures of the SPpellet samples were identified using Fourier transform infrared (FTIR) absorption spectroscopy . Detailed characterization of the chemical bonds was performed using an HR micro-Raman spectrometer  equipped with a 514 nm argon-ion laser. To investigate the changes after charging or discharging, the Li/S cells were disassembled in an argon-filled glovebox and washed three times with diethyl carbonate (DEC) for 5 min to remove lithium salts.2.3.6 in ethylene carbonate (EC) and DEC . Li/S cells were assembled by sequentially stacking the lithium anode, electrolyte, Celgard 2400 separator, and sulfur cathode. Galvanostatic tests of SP450 were performed at room temperature between 1.0 and 4.0 V vs. Li/Li+ at a current density of 15 mA g−1 based on the electrode weight using a battery cycler . The other cathodes were evaluated between 1.0 and 3.0 V. Cyclic voltammetry (CV) was performed at a scan rate of 1 mV s−1 in a voltage rage of 1–3 V. Electrochemical impedance spectroscopy  was collected with an amplitude of 10 mV over a frequency range from 10 mHz to 1 MHz for original and cycled cells.We examined the free-standing SPpellet, SP450compact, and SPconventional samples as sulfur cathodes. Lithium foil  and a polypropylene film (Celgard 2400) were used as the anode and separator, respectively. The electrolyte consisted of 1 M LiPF3.−3, which is lower than the theoretical density of 1.91 g cm−3  and PAN (1.15 g cm−3)). This difference could be related to the presence of pores in the pellet.47 Initially, the apparent density gradually decreased as the temperature increased and then rapidly decreased from 0.9 g cm−3 for SP300 to 0.49 g cm−3 for SP600. In addition, the sulfur content of SP000 was similar to that of the precursor mixture (sulfur/PAN ratio of 80 : 20). However, the sulfur contents of the SPpellet samples decreased with increasing heating temperature, with a rapid decrease observed from 65 wt% for SP300 to 30 wt% for SP600. The reduction in the apparent density and sulfur content was due to the sublimation of sulfur. The XRD patterns of the SPpellet samples are shown in Although considerable research has already been conducted on SPAN powders, this is the first study of bulk-type SPAN electrodes. The electrode synthesis conditions were optimized by investigating the properties of SPAN electrodes prepared at various heating temperatures . Fig. 1a PBM data was replaced with SVG by xgml2pxml:<glyph-data id="z.tbd" format="PBM" resolution="300" x-size="17" y-size="11" xml:space="preserve">0000000000000000000000000000000000000000000000000000111111111111111000000000000000000011111111111111100000000000000000001111111111111110000000000000000000000000000000000000000000000000000</glyph-data><svg xmlns="http://www.w3.org/2000/svg" version="1.0" width="23.636364pt" height="16.000000pt" viewBox="0 0 23.636364 16.000000" preserveAspectRatio="xMidYMid meet"><metadata>Created by potrace 1.16, written by Peter Selinger 2001-2019</metadata><g transform="translate scale" fill="currentColor" stroke="none"><path d="M80 600 l0 -40 600 0 600 0 0 40 0 40 -600 0 -600 0 0 -40z M80 440 l0 -40 600 0 600 0 0 40 0 40 -600 0 -600 0 0 -40z M80 280 l0 -40 600 0 600 0 0 40 0 40 -600 0 -600 0 0 -40z"/></g></svg>N, CH2, and CH bonds, which probably originated from PAN ((C3H3N)n). The spectrum of SP150 was similar to that of SP000, indicating that no chemical transformation of either PAN or sulfur occurred at 150 °C. In contrast, in the FTIR spectra of the samples prepared above 300 °C, the peaks corresponding to CN and CH2 bonds disappeared, implying that both cyclization and dehydrogenation started at 300 °C. The cyclization reaction involves the nitrile groups of the PAN precursor. The FTIR spectra of both SP450 and SP600 contained peaks corresponding to C PBM data was replaced with SVG by xgml2pxml:<glyph-data id="z.dbd" format="PBM" resolution="300" x-size="8" y-size="10" xml:space="preserve">00000000000000000000000000000000111111110000000011111111000000000000000000000000</glyph-data><svg xmlns="http://www.w3.org/2000/svg" version="1.0" width="13.200000pt" height="16.000000pt" viewBox="0 0 13.200000 16.000000" preserveAspectRatio="xMidYMid meet"><metadata>Created by potrace 1.16, written by Peter Selinger 2001-2019</metadata><g transform="translate scale" fill="currentColor" stroke="none"><path d="M0 440 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z M0 280 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z"/></g></svg>C and CN bonds and a hexahydric ring, which are related to the formation of π-conjugated coplanar hexahydric ring structures. As shown in −2, which are characteristic of elemental sulfur. The Raman peaks for SP300 appeared at 1325 and 1530 cm−1, corresponding to the disorder-induced D band and graphitic G band, respectively. This observation indicates that PAN underwent a structural change to form a carbon structure through dehydrogenation and carbonization. In addition, the Raman spectrum of SP300 contained several small peaks at 475, 387, 305, and 177 cm−1, which are characteristic of S–S and C–S bonds. The characteristic peaks in the Raman spectra of SP450 and SP600 were similar to but weaker than those in the spectrum of SP300. The sulfur states of the SPpellet samples were also investigated using TGA, as shown in Fig. S1.8. The TGA curves of SP150 and SP300 were similar to that of SP000. In contrast, SP450 and SP600 were thermally stable to 600 °C. This increased stability is due to the standard bond energy of C–S bonds (740 kJ mol−1) being higher than that of S–S bonds (418 kJ mol−1).48The chemical bonds in the SPpellet samples were further characterized using FTIR and Raman spectroscopy. As shown by the FTIR spectra in 8 and PAN ((C3H3N)n), respectively. In addition, PAN was surrounded by sulfur because SP000 contained more sulfur than PAN. SP150 had fewer bright regions than SP000 and black regions (red arrow) formed on the SP150 surface, which corresponded to pores formed by sulfur melting at 115 °C. The EDS results showed that the bright gray regions (blue arrow) of SP300 were composed of S, C, and N, indicating the presence of SPAN. The area of SPAN in SP300 was similar to that of PAN in SP150, suggesting the transformation of PAN to SPAN. A bright region corresponding to sulfur remained as a film on the surface near the gray region corresponding to SPAN, which is consistent with the XRD and TGA results  and it had an appropriate amount of pores. However, the synthesis of SPAN particles has been reported at temperatures below 300 °C.36 This difference in the optimized heating temperature was due to the electrode structure. Compared with powders, a higher temperature is required for the synthesis of bulk-type SPAN because it is difficult to sublimate sulfur owing to its density. Based on the properties of the SPpellet samples ,57,58 which might be due the SP450 electrode SP450 having a higher porosity than previously reported SPAN-based electrodes. After the 1st discharge, the electrical conductivity of SP450 increased dramatically to 4.66 × 10−4 S cm−1. Considering the difference in electrical conductivity between the previously reported electrodes and the initial SP450 electrode, discharged SP450 is estimated to have a higher electrical conductivity of ∼10−2 S cm−1. Furthermore, the electric conductivity was not restored to its initial value after the 1st charge and 2nd discharge. After cycling, SP450 exhibits semiconductor-like electrical conductivity, which is why the cycling performance is good, even without conducting agents or current collectors.The calculated electrical conductivity of SP450 after cycling is shown in + diffusion into the active materials. The charge-transfer resistance (Rct) for the pristine electrode was ∼450 Ω, which decreased to 280 Ω after the 1st discharge. However, Rct dramatically increased to 2500 Ω after 1st charge and then again decreased to 470 Ω after the 2nd discharge. The initial decrease in Rct may originate from the increase in the electrical conductivity of SP450 after the 1st discharge. The subsequent increase in Rct may be due to increased resistance during charging. Thus, lithiation can decrease the electrical conductivity and Rct.Electrochemical impedance data before and after cycling is shown in Ep, W h kg−1) was calculated using Wc (kg) is the weight of the cell components , V (V) is the potential, and Ccell (A h) is the discharge capacity of the cell. When all the components were considered, the Li/SP450 cell offered a high practical density of 250 W h kg−1 because it did not contain any inactive components and the amount of electrolyte was reduced by using a carbonate-based electrolyte. Specifically, 3.7 μL of electrolyte was used per 1 mg of sulfur, which is similar to the previously reported minimum electrolyte amount. However, the sulfur utilization was two times higher than that in previous reports at similar current densities . Previous studies on the use of SPAN have focused on the cycle performance rather than the electrolyte amount. Typically, for SPAN electrodes, 60 μL of electrolyte has been used per 1 mg of sulfur. If a SP450-based cell was constructed using the same products, such as lithium foil and a separator, the capacity would be 96 W h kg−1, despite SP450 having a higher reversible gravimetric capacity than all reported SPAN-based electrodes. Thus, SP450 exhibited a high practical specific energy density because the free-standing electrode was composed of only active components and a small amount of electrolyte.The practical specific energy density is a major factor affecting the performance of Li/S batteries, with the integration of inactive materials decreasing the practical energy density. Critically, the amount of electrolyte can only be reduced to a limited extent owing to the solubility of the intermediates. The practical Li/S energy density (electrode−1 (1258 mA h gsulfur−1) at the 75th cycle. nd cycle. The average operating potential of SP450-C was higher than that of SP450, which may be due to the addition of conducting agents. The Li/SP450-C cell exhibited an initial discharge capacity of 600 mA h gelectrode−1 (1715 mA h gsulfur−1), which decreased to 470 mA h gelectrode−1 (1343 mA h gsulfur−1) at the 1st charge. Irreversible behavior was also observed, despite the addition of the conducting agents. Both conducting agents and residual Li increase the electrical conductivity of SPAN, although their effects are different.To improve the cell performance, we introduced a small amount of conducting agent (10 wt% MWCNTs) into the electrode because fast electron transfer was difficult in the thick SPAN electrode. The obtained SP450-C electrode had a sulfur content of 35 wt%. As shown in −1. The basic electrochemical characteristics of the cyclic voltammograms are similar to those of the charge–discharge curve profiles. In the 1st cycle, SP450-C exhibited one cathodic peak at 1.2 V. However, in subsequent scans, the onset potential of this peak shifted to a more positive value with broad peaks observed at 1.7 and 2.1 V. These broad peaks were reversible, corresponding to the formation of Li2S from free sulfur after cleavage of C–S and S–S bonds. In addition, the reduction peak initially appeared at 2.4 V and then shifted toward higher potentials during cycling, which suggests that the electrical conductivity improved during cycling, as shown in sulfur−1 at current rates of 0.1, 0.2, 0.5, and 1 C, respectively, which are high values considering the absence of a current collector. This reports shows higher utilization of sulfur with low E/S ratio than previous reports, which shown in Table S2.4.−2 and exhibited a high discharge capacity of 674 mA h gelectrode−1 with a reversible capacity of 517 mA h gelectrode−1 (1260 mA h gsulfur−1) after the 10th cycle. Considering the weight of the anode, separator, cathode, and electrolyte without packaging, the corresponding Li/S cell exhibited a high practical energy density of 250 W h kg−1 because of the small amount of electrolyte (3.7 mL gsulfur−1) and the high gravimetric capacity of SP450. However, to achieve commercialization, the power density of the Li/SP450 cell should be further improved.A free-standing sulfur cathode was prepared by a one-step process, in which compacted sulfur and PAN powders were heated to 450 °C. The SP450 cathode consisted of SPAN without any inactive components, such as a conducting agent, binder, or current collector. The electronic conductivity of SPAN increased more than 15 000 times following lithiation, which may be related to the absence of a conducting agent in SP450. In addition, the SP450 cathode had a high areal loading of 16.37 mg cmThere are no conflicts to declare.RA-011-D1RA02462K-s001


Several glass–ceramic (gc) LPSswith different compositions and good Li conductivity have been previouslyreported, but the relationship among composition, atomic structure,stability, and Li conductivity remains unclear due to the challengesin characterizing noncrystalline phases in experiments or simulations.Here, we mapped the LPS phase diagram by combining first-principlesand artificial intelligence (AI) methods, integrating density functionaltheory, artificial neural network potentials, genetic-algorithm sampling,and ab initio molecular dynamics simulations. Bymeans of an unsupervised structure-similarity analysis, the glassy/ceramicphases were correlated with the local structural motifs in the knownLPS crystal structures, showing that the energetically most favorableLi environment varies with the composition. Based on the discoveredtrends in the LPS phase diagram, we propose a candidate solid-stateelectrolyte composition, (Li2S)x(P2S5)x1– (x ∼ 0.725), that exhibits high ionic conductivity(>10–2 S cm–1) in our simulations,thereby demonstrating a general design strategy for amorphous or glassy/ceramicsolid electrolytes with enhanced conductivity and stability.Lithium thiophosphates (LPSs) with the composition (Li In 2006, Mizuno and co-workers observedthat the conductivity of LPS materials can be further promoted bypartial crystallization of the Li2S–P2S5 glasses.12 By now, a number ofdifferent glass–ceramic (gc) LPS compositionshave been synthesized and characterized, including LiPS3 ((Li2S)0.5(P2S5)0.5),13 Li2PS3 ((Li2S)0.667(P2S5)0.333),16 Li7P3S11 ((Li2S)0.7(P2S5)0.3),25 Li3PS4 ((Li2S)0.75(P2S5)0.25),34 and Li7PS6 ((Li2S)0.875(P2S5)0.125),35 all of which lie on or near the P2S5–Li2S composition line in the Li–P–Sphase diagram (Solid-state batteries (SSBs) are a prospectivealternative to conventionalLi-ion batteries (LIBs), in which the flammable liquid electrolytesare replaced with safer solid Li-ion conductors. Additionally, SSBscan potentially enable the use of Li metal anodes and thus significantlyhigher energy densities. diagram 1.37 spectroscopy as well as with computational methods.41 Nevertheless, glass–ceramic (gc) LPS-basedSEs exhibit both crystalline and noncrystalline phases, and the ionicconductivity of such gc-LPS materials is significantlyinfluenced by the glassy phases.42 Although thecrystal structures and electronic properties of LPS have been thoroughlystudied, the relationship between structures and Li conductivity inthe gc-LPS materials has not been well understood,also due to the limitations of experimental and computational techniquesfor characterizing noncrystalline phases.LPS compositions crystallize in several differentcrystal structures2 that ha44 where different local P–S motifsare present depending on the LPS composition. xSyn– anionicspecies commonly observed: ortho-thiophosphate (PS43–), pyro-thiophosphate (P2S74–), hypo-thiodiphosphate (P2S64–), meta-thiodiphosphate (P2S62–), and meta-thiophosphate (PS3–).45 Polymeric chains ofPS3– are only observed in the LPS glasseswith low Li2S contents modeling and neutron diffraction, it was proposed that the corner-sharingP2S74– shields the positivecharge of P due to electron transfer between P and bridging S, thereforesuppressing Li conduction.67 However, a later ab initio molecular dynamics (AIMD) study found that the flexibility of P2S74– ditetrahedra facilitatesLi+ diffusion.21Previous computational studies mainly focusedon the crystallineLPS phases, such as Li71In essence,only few theoretical studies of amorphous/glassy LPSstructures have been reported, and the effect of amorphization onLi conduction has not yet been well understood. Conventional densityfunctional theory (DFT) based AIMD simulations alone are limited torelatively small structure models with ∼200 atoms, which makesit challenging to investigate amorphous phases without long-rangeordering. In addition, sampling amorphous phases with such moderatelysized structure models using AIMD simulations already required significantcomputational resources. On the other hand, machine learning potentialstrained on first-principles reference data can be efficient and accuratefor describing amorphous phases with reasonable computation cost.gc-LPS with varying composition, we mapped the gc-LPSphase diagram by integrating DFT,72 artificialneural network (ANN) potentials,73 evolutionary/genetic-algorithm(GA) sampling, and AIMD simulations as illustrated by the workflowdiagram in 2S–P2S5 composition line using an AI-aided sampling approach, the phase diagram of gc-LPS was completed. For each LPS composition, GA global structureoptimizations with an ANN potential were performed to determine low-energyatomic configurations. The relevant near-ground-state structures determinedby this sampling approach were recomputed with DFT, and all reportedfinal results are based on DFT. The thermodynamic stability and ionicconductivity of glassy/ceramic phases was correlated with local structuralmotifs by determining similarities of Li sites in glassy and crystallineLPS structures motivated by the recent findings by Sadowski and Albe,64 which allowed identifying structure–composition–conductivityrelationships. With machine learning accelerated sampling and AIMDsimulations, a candidate solid-state electrolyte composition, (Li2S)x(P2S5)x1– (x = 0.724),with high ionic conductivity (>10–2 S cm–1) was identified, which points toward a design strategyfor LPS-basedSE materials with enhanced conductivity and stability.To determine the local atomic structures of 75 and the Perdew–Burke–Ernzerhof (PBE) exchange-correlationfunctional76 as implemented in the Vienna Ab Initio Simulation Package (VASP)74 and an energy cutoff of 520 eV. Gaussian smearing with a width of0.05 eV was used, and total energies were generally converged betterthan 10–5 eV/atom; the final force on each atomwas less than 0.02 eV/Å. The first Brillouin zone was sampledusing VASP’s fully automatic k-point scheme with length parameter Rk = 25 Å.All DFT calculationswere carried out with the projector-augmented-wave (PAW) method77 To obtain near-ground-statestructures as reference for the machine-learning potential (see below),150 evenly spaced snapshots were extracted from the AIMD trajectoriesthat were reoptimized with DFT at zero Kelvin via geometry optimizationsas described above.Amorphousstructure models were generated with AIMD simulations of supercellscontaining 80–128 atoms. In AIMD simulations, a Gamma k-pointscheme was employed to reduce the computational cost. The time stepfor the integration of the equations of motion was set to 1 fs, andthe temperature of the system was set to 1200 K using a Nosé–Hooverthermostat.69To determine ionic conductivities, ∼300ps long AIMD simulationswere performed for select compositions  centered on atom i is approximated asmax is the expansion order, ϕα and ϕ̅α are the basis function corresponding to order α as wellas its orthogonal dual function, and cαRDF are theexpansion coefficients. The second sum in the expression of the coefficientsruns over the Cartesian coordinates R⃗j of all atoms j within thelocal environment of atom i, σi, the distance between atom i and its neighbor j is denoted Rij, and fc is a cosine cutoff function that smoothlygoes to zero at a defined maximal interaction distance. The expansionof the angular distribution function (ADF) is equivalent and yieldsthe expansion coefficients {cαADF}.To be suitableas inputs for our machine-learning models, local atomic environments,including atomic positions and species, need to be featurized, i.e.,transformed to a vector representation with constant dimension.cαRDF} and {cαADF} are invariantwith respect to the rotation and translation of the atomic structureand the permutation of equivalent atoms, which makes them suitablefeatures of the local structure. To incorporate information aboutthe chemical species within the local atomic environment σi, the contribution of each atom j is weighted with an element-specific weight wtj (tj is the type of atom j), yielding a second set of expansion coefficients {c̃αRDF} and {c̃αADF}. The complete feature vector of thelocal atomic environment of atom i is then givenby the concatenation of the four sets of expansion coefficients79 Hence, the dimension of the Chebyshevfeature vectors f⃗i is 2 × (19 + 5) = 48, including also the coefficients for expansionorder 0. We used wtj = −1, 0, +1 to weight the contributionsof the three species.The RDF and ADFexpansion coefficients {80 as implemented in the atomicenergy network package (ænet).81 ANN potentials represent the total energy Etot of an atomic structure as the sum of atomicenergies, Etot = ∑iNatomEi, where the atomic energies Ei are predicted by ANNs for a given local atomic environmentand Natom is the number of atoms in thestructure. Local atomic environments were represented as describedabove. An ANN architecture with two hidden layers of 15 nodes eachand hyperbolic tangent activation functions was employed. The Broyden–Fletcher–Goldfarb–Shanno(BFGS) method82 was employed for the weightoptimization. A total of 10% of the reference data were randomly selectedas an independent validation set for cross-validation and were notused during training. The training was repeated 10 times for 500 trainingiterations using different randomly initialized weight parameters,and the ANN potential with the lowest validation-set error was selected.All machine-learning potential(MLP) simulations were performed with artificial neural network (ANN)potentialsgc-LPS structure configurations were generated with the genetic algorithmsampling approach described below using the ANN potential; and (iii)the 10 structures with lowest ANN potential energy among those sampledwere reoptimized using DFT and added to the reference data set. Thefinal ANN potential yields a root-mean-squared error of 1.4 meV/atomand a mean absolute error of 0.6 meV/atom relative to the DFT referenceenergies in an independent validation set that was not used for trainingand contained 10% of the structures in our database. As previouslydemonstrated for amorphous LiSi alloys and LiPON solid electrolytes,70 specialized ANN potentials constructed based on moderately sizedreference data sets can be used in conjunction with DFT for acceleratedsampling of amorphous phases.For accelerated sampling, a specialized ANN potential was trainedon a data set containing ∼6000 atomic structures that weregenerated with the following iterative approach: (i) An initial ANNpotential was trained on the LPS crystal structures with lattice parametersscaled between ±15% and randomly perturbed atomic positions fromshort AIMS simulations at 1200 K; (ii) a number of 2S–P2S5 composition line were sampled with a genetic-algorithm(GA) as implemented in the atomistic evolution (ævo) package(http://ga.atomistic.net),68 following previously reported strategiesthat are briefly described in the following.70 Although glassyphases lack long-range ordering, it can be expected that the localatomic motifs in gc-LPS phases resemble those ofthe known LPS crystalline phases (x in (Li2S)x(P2S5)x1– via removing Li2S or P2S5, respectively, from supercellsof the known LPS crystal structures. The approach is as follows:1.3, Li7P3S11, β-Li3PS4, γ-Li3PS4, or Li7PS6 is chosen as the parent structure;A supercell of one of the crystal structuresLiPS2.2S and P2S5;The GA is used toidentify combinationsof 2 Li and 1 S atoms that can be removed with low formation energyrelative to Li3.2S deficientcomposition is optimized with DFT; andThe created Li4.The optimized structure is taken tobe the new parent structure, and the algorithm continues with step(2).With the specialized ANNpotential, the amorphous phases along the Lie phases 3. The ph2S–P2S5 composition line by removing 2 P and 5 S atoms at each step (insteadof 2 Li and 1 S atoms).We used the same technique to sample in the oppositedirection on the LiThe GA employed a population size of32 trials and a mutation rateof 10%. For each composition, at least 10 lowest energy structuremodels identified with the ANN-GA approach were selected and fullyrelaxed with DFT to obtain the first-principles phase diagram. Weemphasize that the GA sampling approach yields, by design, DFT optimizedstructures and their DFT energies.2S)x(P2S5)–x1 the correspondingformation energy per atom was calculated asE is the total energyof a specific configuration as predicted by DFT; x is the molar fraction of Li2S in the LPS composition;and E2SLi and E2S5P are constant and are equalto the total energy per formula unit of bulk Li2S and P2S5, respectively. For any given composition, theconfiguration with a lower formation energy is thermodynamically favoredat zero Kelvin. The stabilities of different compositions can be comparedby constructing the lower convex hull of the formation energies toobtain the phase diagram.68For any given structure and composition is the mean atomicfeature vector (the first moment). The structure fingerprint is thenthe union  of the distribution moments, F⃗ = f⃗(1) ∪ f⃗(2) ∪ ..., until a maximum moment.In practice, we found that truncating after the second moment alreadyyielded unique structure fingerprints that can distinguish all atomicstructures in our database. Atom-type specific structure fingerprintscan be constructed by including only atomic feature vectors for thelocal atomic environments of select atomic species. We made use ofthis approach by constructing structure fingerprints based on onlythe local atomic environment of Li atoms. Finally, we reduced thedimension of the structure fingerprints by performing a principalcomponent analysis (PCA) after data standardization, using the PCAand StandardScaler implementations of the scikit-learn library.83 We found 10 principal componentsto be sufficient, which can explain 85% of the data variance. Hence,each atomic structure in our database could be uniquely representedby a fingerprint vector with 10 components.Low-energyamorphous LPS structures were compared with the known LPS crystalstructures by their connectivity of PSSp of two atomic structures as the Pearson correlation coefficientF⃗1 and F⃗2 are two (dimension-reduced)structure fingerprints. Furthermore, we performed a cluster analysisof the structure fingerprints using the k-means approachas also implemented in scikit-learn.83Using the structurefingerprints, we define the similarity 2S–P2S5 compositionline started with 13 LPS crystal structures with the formula unitsLiPS3,13 Li2PS3,16 Li4P2S7,60 Li7P3S11,17 α-Li3PS4,31 β-Li3PS4,32 γ-Li3PS4,30 Li48P16S61,84 and low-temperature(LT)-Li7PS635 thathad previously been reported based on experimental characterizationand/or theoretical modeling. The crystal structures, which were obtainedfrom the Inorganic Crystal Structure Database (ICSD)85 and the Materials Project (MP)86 database, are shown in 2S and P2S5,the end points of the composition line, are shown in 3PS4)appears on the lower convex hull of the formation energies and isthus predicted to be thermodynamically stable at zero Kelvin. Thepreviously reported superionic conductors, β-Li3PS432 and Li7P3S11,17 are 3.2 meV/atom and 17.2 meV/atomabove the convex hull, indicating that they are metastable at zeroKelvin. However, the energy difference between β-Li3PS4 and γ-Li3PS4 is small(3.2 meV/atom) compared to the thermal energy per degree of freedomat room temperature (∼26 meV), so that it is plausible thatthe β polymorph can be thermodynamically stable at room temperature.Note that the crystal structure of Li4P2S760 is a theoretical prediction from the literatureand has not been characterized experimentally yet, which is consistentwith its comparatively high decomposition energy of 23.5 meV/atomin our phase diagram.Our computational samplingof the Li2S or P2S5 from supercells of the crystal structures. This compositionsampling yielded low-energy structures with structural disorder andno symmetry, as one would expect for amorphous or glassy phases, while still exhibiting local similaritieswith the parent crystal structures from which they were derived. Atzero Kelvin, these glass–ceramic structures are also predictedto be thermodynamically unstable, though they might be stabilizedat synthesis temperatures due to their high entropy (entropy control)or via kinetic trapping.Also shown in the phase diagram of 3 and Li4P2S7 and between Li3PS4 andLi7PS6, respectively. This means that compositions(Li2S)x(P2S5)–x1 with 0.5 < x < 0.667 and 0.75 < x < 0.875will likely phase separate instead of forming a solid solution, inagreement with previous experimental observations  were found. It can, therefore, beexpected that compositions with 0.667 < x <0.75 can be more readily synthesized.As seen in the phase diagram, the ANN-GAsampling identified twomiscibility gaps between LiPS45 Isolated PS43– tetrahedraare mostly observed in the gc-LPS compositions withhigh Li2S content (x ≥ 0.75), suchas α-Li3PS4, β-Li3PS4,32 γ-Li3PS4,30 and Li7PS6.35 The P2S74– motif, consisting of two corner-sharing PS4 tetrahedra,is the main building block of the Li7P3S11 crystal structure17 as well asglassy LPS compositions with x < 0.75. The P2S62– motif, formed by two edge-sharingPS4 tetrahedra, is observed in gc-LPSwith x ≤ 0.6 and is the only local motif inLiPS3 crystals.13 The P2S64– with direct P–P bondingis typically present in gc-LPS with 0.6 ≤ x ≤ 0.7.24 Note that theoxidation state of P is +4 only in the P2S64– motif, while it is +5 in all other local motifs.The P2S64– motif also occursin Li2PS3,16 which is a sulfur-deficientcomposition that is not on the Li2S–P2S5 composition line.The LPScrystal structures shown in gc-LPS phases, wecomputed the radial pair distribution functions (RDFs) for P–Sand Li–S in gc-LPS compositions with 0.385≤ x ≤ 0.867 as shown in Figure S1. As seen in the figures, and as expected, the RDFs of thegenerated gc-LPS structures exhibit features of thecrystal structure RDFs but show broadened peaks with shifted peakpositions. In general, with decreasing amount of Li2S in gc-LPS, the main Li–S peak shifts to greater distances,which is caused by the formation of corner-sharing motifs, in agreementwith previous reports.61 Note that for a large fraction of the gc-LPS structures(∼1/3) the shape of the RDF differs significantly from thatof the parent structure; i.e., the RDFs of derived structures exhibitdifferent peaks than the RDF of the parent crystal structure. Instead,structures with the same composition that were derived from two differentparent structures exhibit similar peaks, indicating that these compositionshave a strong preference for specific structural motifs. This is especiallyevident in the S–S RDF shown in Figure S1 and indicates that gc-LPS with compositionsin between the crystalline phases may exhibit multiple different localstructural motifs found in the neighboring (by composition) crystallineLPS.To better understand the localstructures of the ANN-GA sampled As discussed in the 3, Li7P3S11, β-Li3PS4, γ-Li3PS4, Li7PS6, and Li2S are shown. The Pearsoncorrelation Sp of thestructure descriptors . Most of the structures derived fromeither LiPS3 or Li7PS6 remain similarto their parent structure during sampling, leading to distinct clustersfor these structures in 3PS4 (x = 0.75). Within thenarrow composition range 0.70 ≤ x ≤0.75, the structures closest to the ground-state hull change in characterfrom Li7P3S11 to structures thatare similar to β-Li3PS4 and γ-Li3PS4.In tors see  sectionture see S2. Most k-means clustering.The predicted grouping resembles the one shown in 3PS4, the cluster analysis finds that the Li environment changes withincreasing energy, which we can attribute to the γ, β,and α polymorphs. At high energies above the ground state hull,a fourth class of the Li environment is found of which Li7P3S11 is also a member, though it is unlikelythat these structures can be synthesized at any conditions.Instead of classifying the sampled glass–ceramicstructuresby their similarities to reference crystal structures, gc-LPS phases with compositions betweenLi7P3S11 (x = 0.70)and Li3PS4 (x = 0.75) exhibitthe same type of Li environments as the superionic conductor β-Li3PS4. Given this energetic preference, it is likelythat β-Li3PS4-like Li environments arepresent in as-synthesized gc-LPS within this compositionrange or would form over time. To determine if this similarity alsotranslates to Li conductivity, we performed AIMD simulations for aglass–ceramic LPS with composition gc-Li42P16S61 (x = 0.724),the two neighboring crystalline phases (β-Li3PS4 and Li7P3S11), and a compositionoutside the target range, gc-Li38P24S79 (x = 0.613), for comparison.The ionic conductivities at room temperature were obtained from Arrheniusextrapolation  and are compiled in gc-LPS from theliterature, which are sensitive with respect to the experimental conditions,e.g., temperature and pressure. Samples prepared under different conditionsmay exhibit different local motifs, leading to a wide range of measuredconductivities.23The cluster analysis of the Li atomenvironments discussed in the previous section indicates that thelowest-energy 7P3S11 is in good agreement withpreviously reported experimental measurements and theoretical calculations.The differences are greater for the β-Li3PS4 phase, where the agreement with previous simulations is good butpredicted conductivities are significantly greater than those observedin experiments. This has to be expected, since the metastable βphase is more challenging to characterize experimentally as well asin simulations. Hence, the data for the β phase is subject togreater uncertainties.As shown in gc-Li42P16S61 is high (33.1 mS cm–1) andlies between the conductivities of crystalline Li7P3S11, 46.9 mS cm–1, and β-Li3PS4, 14.3 mS cm–1. In comparison,the other amorphous phase, gc-Li38P24S79 (x = 0.613), has a significantlylower ionic conductivity of 3.45 mS cm–1 and higheractivation energy of 0.282 eV  in the phase diagramx(P2S5)x1– with 0.5 ≤ x ≤ 0.667 and 0.75 ≤ x ≤0.875, predicting that solid solutions with such compositions wouldbe challenging to synthesize and likely to phase separate at roomtemperature. Dietrich et al. previously conducted an experimentalstudy of glass–ceramic LPS compounds with 0.6 ≤ x ≤ 0.8 and found that LPS (x =0.8) phase separates into Li3PS4 (x = 0.75) and Li2S (x = 1.0),24 in agreement with our prediction. However, thesame authors reported the successful preparation and characterizationof LPS (x = 0.6), which should also be unstable basedon the calculated phase diagram. A possible explanation for this discrepancycould be sulfur deficiency in the compositions, since Li4P2S6 is a known decomposition product of gc-Li4P2S713 and an attractor in the phase diagram  to x < 0.75. This relative destabilizationof the γ phase is visualized in gc-Li42P16S61 (x = 0.724) exhibits a highLi conductivity of 33 mS cm–1. The RDF analysisof gc-Li42P16S61 derived fromβ-Li3PS4 still resembles that of the parentphase. As seen in gc-Li42P16S61 structure exhibits both well-ordered and disordered domains,and the Li probability distribution is greater in the ordered regions.This further indicates that reminiscence of the crystalline phaseis important for Li conductivity in this gc-LPS composition.Though we note that the PSx motifs donot generally control the Li environments, there are structures withsimilar P–S RDFs but different Li environments. An exampleis analyzed in Supporting Information Figure S4.The calculatedphase diagram shows that the superionic LPS compoundsare metastable and therefore prone to decomposition, which is in agreementwith previous experimental and computational work discussed in the ronments 7 points 2S5 and Li3PS4 at zero Kelvin; i.e., it is abovethe convex hull of formation energies. It has previously been establishedthat knowledge of the energy above the convex hull is insufficientto predict synthesizability87 and thatthe thermodynamic limit for the synthesis of metastable compoundsis chemistry-dependent.88 On the otherhand, a wide range of different gc-LPS compositionshave previously been reported x(P2S5)x1– with x = 0.724, which means that the phase, if it can be synthesized, couldbe expected to be shelf-stable at room temperature.Similar to the known crystalline LPS superionicconductors, thehere identified LPS composition is also metastable and thermodynamicallyunstable with respect to decomposition into P see ref ), indica2S–P2S5) are known butare unstable due to phase transitions, the local atomic environmentof the Li sites can be taken as a design target, in agreement withprevious findings.64 (2) Potentially stablesuperionic conductors can then be identified by searching for regionswithin the composition space that energetically favor the target Lisite environment over other environments.Taken together,the observations made in the present work led tothe following design strategy for amorphous solid Li conductors: (1)If Li superionic conductors within a given composition space x(P2S5)x1–, solidelectrolytes using first-principles calculationswith AI-aided sampling and structure similarity analysis. The phasediagram exhibits two pronounced miscibility gaps, so that compositionswith 0.5 < x < 0.667 and 0.75 < x < 0.875 are prone to phase separation at room temperatureeven if they can be synthesized. We showed that glassy/ceramic phaseswith compositions 0.70 < x < 0.75 are morelikely to be stable because of their lower decomposition energiesand exhibit Li sites with local structural environments similar tothose in the superionic conductor β-Li3PS4. This led us to propose a candidate solid-state electrolyte composition,(Li2S)x(P2S5)x1–, with x = 0.724, that exhibits high ionic conductivity (>10–2 S cm–1) in simulations, demonstrating a designstrategy for glassy or amorphous solid-electrolyte materials withgood conductivity and stability.We mapped the phase diagram of lithiumthiophosphate, (Li