Abstract
Solid-state Li–S batteries (SSLSBs) are made of low-cost and abundant materials free of supply chain concerns. Owing to their high theoretical energy densities, they are highly desirable for electric vehicles1,2,3. However, the development of SSLSBs has been historically plagued by the insulating nature of sulfur4,5 and the poor interfacial contacts induced by its large volume change during cycling6,7, impeding charge transfer among different solid components. Here we report an S9.3I molecular crystal with I2 inserted in the crystalline sulfur structure, which shows a semiconductor-level electrical conductivity (approximately 5.9 × 10−7 S cm−1) at 25 °C; an 11-order-of-magnitude increase over sulfur itself. Iodine introduces new states into the band gap of sulfur and promotes the formation of reactive polysulfides during electrochemical cycling. Further, the material features a low melting point of around 65 °C, which enables repairing of damaged interfaces due to cycling by periodical remelting of the cathode material. As a result, an Li–S9.3I battery demonstrates 400 stable cycles with a specific capacity retention of 87%. The design of this conductive, low-melting-point sulfur iodide material represents a substantial advancement in the chemistry of sulfur materials, and opens the door to the practical realization of SSLSBs.
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The data that support the findings of this study are available from the corresponding author upon request.
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Acknowledgements
This work was supported by the Advanced Research Projects Agency–Energy, US Department of Energy (DOE), under contract no. DE-AR0000781. M.L.H.C. and S.P.O. acknowledge the support from the Materials Project, funded by the US DOE, Office of Science, Office of Basic Energy Sciences, Materials Sciences and Engineering Division, under contract no. DEAC02-05-CH11231 (Materials Project Program No. KC23MP). Computing resources were provided by the National Energy Research Scientific Computing Center (NERSC) and the Advanced Cyberinfrastructure Coordination Ecosystem: Services & Support (ACCESS) under grant no. DMR-150014. S.T. and E.H. are supported by the Assistant Secretary for Energy Efficiency and Renewable Energy (EERE), Vehicle Technology Office of the US DOE through the Advanced Battery Materials Research (BMR) Program under contract no. DE-SC0012704. This research used 28-ID−2, 8-BM and 7-BM beamlines of the National Synchrotron Light Source II, US DOE Office of Science User Facilities, operated for the DOE Office of Science by Brookhaven National Laboratory under contract no. DE-SC0012704. C. Wu, Z.F. and Y.Y. are supported by the US DOE’s Office of EERE under the Vehicle Technologies Program under contract no. DE-EE0008864. This work made use of the shared facilities of the UC Santa Barbara MRSEC (grant no. DMR-720256), a member of the Materials Research Facilities Network (http://www.mrfn.org). This research used the Electron Microscopy facility of the Center for Functional Nanomaterials (CFN), which is a US DOE Office of Science User Facility, at Brookhaven National Laboratory, under contract no. DE-SC0012704. Canhui Wang and Chao Wang were supported by the Advanced Research Projects Agency–Energy, US DOE, under contract no. DE-AR0001191. FIB and SEM characterizations were performed at the San Diego Nanotechnology Infrastructure (SDNI) of UCSD supported by the NSF (grant no. ECCS1542148). Raman facilities were supported by the NSF through UCSD MRSEC, grant no. DMR-201192. The authors acknowledge the use of facilities and instrumentation at the UC Irvine (UCI) Materials Research Institute (IMRI), which was supported in part by the NSF through the UCI MRSEC (grant no. DMR-2011967). XPS facilities were funded in part by the NSF Major Research Instrumentation. We thank the Molecular Mass Spectrometry Facility at UC San Diego for performing the MALDI-TOFMS measurement, which is supported by the NSF under grant no. CHE−1338173. Part of this work was performed in the Cordx Yufeng Li Collaboratory.
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J.Z. and P.L. conceptualized the idea and designed all of the experiments. S.P.O. and M.L.H.C. performed the theoretical calculations and drafted the results. S.T. and E.H. conducted the XAS and X-ray PDF tests. S.W. performed the Raman, DSC, cryo-SEM and XPS measurements. C. Wu, Z.F. and Y.Y. collected the cross-sectional SEM images of the full cells at different states, including anode interfaces and cathode interfaces. H.N. collected EPR data, H.N. and R.J.C. analysed and wrote up the results. Canhui Wang and Chao Wang performed the cryo-TEM. Y.X. and E.E.F. performed the XRD tests. Q.R.S.M. did the in situ heating XRD test. H.L., S.Y., G.H., J. Holoubek, J. Hong and C.S. helped with the synthesis of materials, cell fabrications, electrochemical performance tests, data analysis, discussion and revision. C.J.B. helped with the revised manuscript preparation and discussion. J.Z., P.L. and S.P.O. drafted the manuscript with input and revision from all authors. P.L. and S.P.O. supervised the research. J.Z. and M.L.H.C. contributed equally to this work.
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P.L. and J.Z. report a US provisional patent application filed on February 13, 2023, Serial No. _63/484,659, based on this work.
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Extended data figures and tables
Extended Data Fig. 1 Healable interfaces in SSLSBs with a low-melting-point sulfur iodide, and the synthesis and characterization of sulfur iodide materials.
(a) Schematic of a solid state battery with elemental sulfur as the active material. Poor solid/solid contact develops during cycling due to volume changes of the active material. (b) Schematic of a solid state battery with sulfur iodide as the active material. Ideal active material/electrolyte interface is achieved through periodical heating to melt the cathode, thus healing the interface. (c) Illustration of the procedures to prepare S9.3I. (d) The DSC curves of sulfur iodide with different S:I ratios in the temperature window of 20−150 °C. With S:I ratios decreasing from 1:0 to 9.3:1, the typical sulfur endothermic peak gradually disappears to transition to one phase melting behavior. Beyond 6:1, the exothermic peaks assigned to iodine are observed. (e) Spectra obtained during in situ heating XRD of S9.3I from 30 to 130 °C with an 5 °C interval. (f) Raman spectra of sulfur iodide with different ratios, which features an iodine peak located at 423.8 cm−1 when the S:I ratio is less than 9.3:1. The cryo-SEM image (g) and the corresponding EDX mapping images of I (h) and S (i) element distribution of S9.3I after melting. (j) XRD of S, I2, S35.7I, S15.9I, S9.3I, S6I. In the low angle range, the relative intensity of the broadened peak centered at ~1.5° gradually increases with the increase of iodine, but from 9.3:1 to 6:1, the iodine diffraction peaks also emerge. (k) Mass spectra of S and S9.3I. Compared to S, S9.3I shows two additional peaks centered at 127.0 and 253.8 m/z, which are attributed to I− and I2 species respectively. No peaks can be indexed to species containing S-I bonds. (l) The PDF of sulfur iodide materials after melting with different S and I ratios, elemental sulfur and iodine. When the ratio of S to I is below 9.3:1, bulk I2 is present as indicated by the signals from the short and long range structures.
Extended Data Fig. 2 Computed Crystal and electronic structure of S-I compounds.
(a) Computed XRD patterns of elemental S and predicted S9.6I structure. (b) Computed PDF of S9.6I. We attribute any minor discrepancies to the fact that the DFT relaxations are carried out at 0 K, while the experimental data is obtained at 300 K. (c) Computed XRD patterns of S, S9.6I, S8I and S with S8 ring vacancy structure. (d) HSE projected density of states of S, S9.6I and S8I. Elemental projected density of states for S (e) and S9.6I (f) using SCAN functionals.
Extended Data Fig. 3 Electronic structure and conductivities of sulfur iodide materials.
(a) Variable temperature X-band EPR spectra collected on S9.3I and elemental sulfur. The radical concentration in S9.3I grows with increasing temperature, but not in S. (b) Temperature evolution of the EPR signal intensity (top panel), linewidth (bottom panel), and g-factor (middle panel) of S9.3I. The g-factor for S9.3I is approximately 1.99 and does not change appreciably with increasing temperature (bottom panel), suggesting that similar radical species are produced over the entire temperature range. (c) EPR spectra collected at 100 K on S9.3I before and after melting. The electronic conductivity of S35.7I (d), S15.9I (e), and S6I (f) measured by using a potentiostatic test at room temperature.
Extended Data Fig. 4 SEM and elemental mapping of S and S9.3I mixed with LPS after heating at 100 °C.
The SEM image of S/LPS mixture (a) and S9.3I/LPS mixture (b) after milling and heating at 100 °C. (c) The cross sectional cryo-FIB SEM images and corresponding element distribution images of S/LPS for S and P. The distribution of S and P is inhomogeneous. (d) Distribution of S, P, and I in S9.3I/LPS. The distribution in the entire area is homogeneous.
Extended Data Fig. 5 Electrochemical performance data of solid state cell components.
The cycling stability and capacity of a LPS/VGCF cathode at 100 °C (a) and 25 °C (c). The corresponding voltage profiles at 100 °C (b) and at 25 °C (d). Capacity contribution of LPS is very limited and the cycling stability of LPS is very poor either at 100 °C or 25 °C with low coulombic efficiencies. The long-term cycling stability of Li/SP/LPSCl/SP/Li symmetric cells at 100 °C (e) and 25 °C (f) at a current density of 0.3 mA cm−2.
Extended Data Fig. 6 The performance of sulfur iodide cathode at different conditions and their comparison with other SSLSBs.
The dQ/dV curves during delithiation of S9.3I (a) and S (b) cathode based on the charging curves at 100 °C from Figs. 2a and b. (c) The cycling stability of S, S35.7I, S15.9I, S9.3I and S6I cathode at 1.6 A g−1 and at 100 °C. (d) The corresponding cycling coulombic efficiency in (c). (e) Rate capability of S9.3I cathode at 100 °C at current densities of 0.32, 0.48, 0.80, 1.28, 1.60, 2.40, 3.20, 4.00, 4.80 and 5.60 A g−1. Rate performance comparison between Li-S9.3I solid state batteries with other SSLSBs: based on the weight of (f) only S and (g) S composite, red-colored symbols represent Li metal anode based SSLSBs while other cells used non-Li metal anodes. (h) The cycling performance of S9.3I cathode with a high mass loading of 4.2 mg cm−2 at 0.32 A g−1 and at 100 °C. (i) The corresponding voltage profiles at different cycles in (h). (j) Cycling stability comparison between Li-S9.3I solid state batteries with other reported SSLSBs. The cycling performance of the S9.3I cathode at varied temperatures and current densities: (k) 80 °C, 0.8 A g−1, (l) 60 °C, 0.8 A g−1, (m) 40 °C, 0.32 A g−1. It should be noted that a low current density was applied to the cells tested at all temperatures during the initial formation cycle.
Extended Data Fig. 7 Morphological study of the Li-SP-LPSCl interface during cycling in full cells.
The cross-sectional SEM of Li/SP/LPSCl interface at different states: pristine (a), fully discharged to 1.5 V (b) and fully charged back to 3.0 V (c). It should be noted that all of the anode interfaces were cut from the full cells with the configuration of S9.3I-LPS-VGCF/LPS/LPSCl/SP/Li with a high mass loading of S9.3I 4.2 mg cm−2 running at 100 °C and at a current density of 0.32 A g−1. At the pristine state (a), both the Li/SP interface and SP/LPSCl SSE interface are well connected. After discharge (b), the thickness of Li decreases from ~50 μm to ~30 μm, the 20 μm difference corresponding to an area capacity of ~4.0 mAh cm−2. After charging back to 3.0 V (c), the thickness of Li recovers back to ~50 μm. Importantly, both the Li/SP interface and SP/LPSCl SSE interface are well maintained without the growth of Li dendrites during the discharge and charge processes.
Extended Data Fig. 8 EIS study of Li-S9.3I and Li-S full cell at 25 °C during initial cycles.
(a) The initial discharge/charge curves of a Li-S9.3I cell, and the corresponding dQ/dV curves (b) derive from (a). (c) The EIS of the Li-S9.3I cell at different states of discharge. (d) The EIS of the Li-S9.3I cell at different states of charge. The DRT analysis of EIS at different states of discharge (e) related to (c), and states of charge (f) related to (d). (g) The initial discharge/charge curves of a Li-S cathode, and the corresponding dQ/dV curves (h) derive from (g). (i) The EIS of the Li-S cell at different states of discharge. (j) The EIS of the Li-S cell at different states of charge. The DRT analysis of EIS at different states of discharge (k) related to (i), and states of charge (l) related to (j). The impedance of S9.3I cathode is much lower than that of pure S cathode during cycling.
Extended Data Fig. 9 Characterizations of intermediates during cycling to diagnose the S9.3I cathode working mechanism.
(a) The corresponding ratios of S0, Sb, St, and S2- in the S9.3I cathode in Fig. 3b at different discharge/charge states. (b) Ex situ Raman spectra of S9.3I during the discharge/charge processes. The ex situ Raman spectra suggest the good reversibility of S9.3I cathode at RT. X-band EPR spectra obtained on (c) the pristine S9.3I cathode, and on (d) the SSE and carbon mixture (LPS-VGCF). (e) Ex situ EPR data obtained on cathode samples harvested from cells stopped at various stages of the initial discharge/charge processes. (f) Parameters obtained from fits of the EPR spectra in (e), including the ratios of the signal intensities from the S9.3I cathode material and from the VGCF additive, and the linewidth of the S9.3I EPR signal. The ratio of S9.3I radicals to carbon decreases during initial discharge and grows upon subsequent charge (f). This trend is consistent with the following model, assuming that polysulfide chains that form in the solid-state exist as dianions and cannot be detected in our EPR measurements, unlike EPR-active polysulfide radicals formed in solution through disproportionation of polysulfide dianions (as in conventional Li-S cells). (g) The S2p high resolution XPS of comparison samples, including LPS/VGCF mixture, S cathode discharge to 1.3 V (D-1.3 V) and chemically lithiated S9.3I prepared by reacting with Li powder with an equivalent capacity of ~800 mAh g−1. (h) Digital photos of elemental S cathode discharged to 1.3 V, S9.3I cathode discharged to 1.3 V, chemically lithiated S9.3I immersed in THF solution and standard Li2S4, Li2S6 THF solutions. (i) The UV-Vis spectra of corresponding samples in (h). The S9.3I cathode discharged to 1.3 V and chemically lithiated S9.3I suspensions exhibit the color of LiSx solutions immediately, but elemental S cathode discharged to 1.3 V doesn’t show any change.
Extended Data Fig. 10 Structure and property evolution of S9.3I cathode during cycling.
(a) XRD of chemically lithiated S9.3I. A chemically lithiated S9.3I particle with a few micrometers (b) and corresponding iodine element mapping (c) shows a uniform iodine distribution. EDX mapping images of S9.3I cathode at different states: (d) Pristine, (e) after 10 cycles and heated over 48 h at 100 °C, and the corresponding element analysis. There is only a very small S: I atomic ratio change after 10 cycles and then heated at 100 °C for over 48 h, indicating the thermal stability of S9.3I cathode. (f) EIS spectra and (g) corresponding DRT analysis of Li-S9.3I full cell at different cycles.
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Zhou, J., Holekevi Chandrappa, M.L., Tan, S. et al. Healable and conductive sulfur iodide for solid-state Li–S batteries. Nature 627, 301–305 (2024). https://doi.org/10.1038/s41586-024-07101-z
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DOI: https://doi.org/10.1038/s41586-024-07101-z
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