Strategies of Enhancing Ionic Conductivity in Model Solid‐State Electrolyte Li15P4S16Cl3

High ionic conductivity is the key point in the development of new solid‐state electrolytes. Herein, a combining strategy of anion (O2−) doping and structure distortion is applied to enhance the Li+ ion conductivity in Li15P4S16Cl3, thus converting the nonionic conductor into fast ionic conductor. Solid‐state 6Li nuclear magnetic resonance analysis shows redistribution of Li+ ions in Li15P4S16Cl3 with O2− doping or local structure distortion via ball milling, indicating energy changes at different lithium sites. As a result, the activation energy is reduced from 0.50 to 0.35 eV for the ball‐milled Li15P4S15.6O0.4Cl3, and the ionic conductivity is enhanced from 10−9 to 10−4 S cm−1. The electrochemical stability of Li15P4S15.6O0.4Cl3 is broadened at the anode side as well. The symmetric cell Li|Li15P4S15.6O0.4Cl3|Li can cycle more than 1000 h with negligible voltage increase. The LiCoO2|Li15P4S15.6O0.4Cl3|Li‐Si all‐solid‐state battery demonstrates an initial capacity of 106 mA h g−1 and retains 92% capacity after 200 cycles at 0.5 C, highlighting excellent rate performance and electrochemical stability.


Introduction
[9] Among them, high ionic conductivity is the key factor limiting the diversity of SSEs.Li þ ionic conduction in solids is achieved through the migration of Li þ ions between adjacent sites as claimed by classical theory.During this process, the site energy of Li þ ions varies with the changes in their surrounding chemical environment in terms of anion type, number, and bond length. [10]The maximum energy difference of Li þ ions during migration is the activation energy of Li þ ion diffusion, which determines the probability of successful Li þ ion transport between adjacent sites.When the stable Li þ sites are facesharing, the migration path is short and the activation energy is relatively low. [11]or example, in cubic Li 7 La 3 Zr 2 O 12 , Li þ ions at 24 d site and 96 h site share crystal faces, enabling fast ion migration in the path of 24 d-96 h-96 h-24 d. [12,13] In bodycentered sulfides, such as Li 10 GeP 2 S 12 , Li 7 P 3 S 11 , and LiZnPS 4 , superionic conductors, [11,[14][15][16][17][18] Li þ ions occupy face-sharing tetragonal sites and migrate fast between adjacent sites (T-T). [19,20]However, according to Pauling's third rule, the presence of face-sharing sites reduces the structure stability. [21]In most cases, the occupied Li þ sites are edge-sharing or core-sharing, in which the ion migration activation energy is relatively high. [22]Like in Li 2 S, stable tetrahedral sites are edgesharing, Li þ ions cannot migrate directly between tetrahedral sites but from tetrahedral sites through interstitial octahedral sites to adjacent tetrahedral sites (T-O-T). [11]In this case, the high site energy at octahedral interstitial sites significantly increases the migration activation energy of Li þ ions (>0.4 eV), making Li 2 S a nonionic conductor. [11]To reduce the migration activation energy and improve ion conductivity, a direct approach is to reduce the energy difference between tetrahedral sites and octahedral sites.In solid materials, Li þ ions are mainly subjected to electrostatic interactions with surrounding ions, especially the nearest neighboring anions.These electrostatic interactions are determined by the type of ion, distance, etc. (surrounding chemical environment).[35] As in LiAlCl 4 , partly Li þ ions migrate to the tetrahedral sites after ball milling, indicating the decreasing of the site energy at tetrahedral sites. [33,34]ere, we take Li octahedral sites) through anion doping and local structure modification, thus reducing Li þ ion migration activation energy and therefore enhancing ionic conductivity.Li 15 P 4 S 16 Cl 3 is a cubic sulfide with Li þ ions distributed in two core-sharing tetrahedral sites (12a and 48e). [22]In this work, by optimized O 2À doping and ball milling, the site energy of Li þ ions at octahedral site decreases, making this site relatively stable and forming a continuous ion migration channel between face-sharing sites (T-O-T).A reduced activation energy of 0.37 eV and enhanced ionic conductivity 10 À4 S cm À1 are achieved, converting Li 15 P 4 S 16 Cl 3 from a nonionic conductor to an ionic conductor.These results indicate that the site energy regulation method based on microstructure optimization can significantly enhance the ion conductivity of solid materials, which may help to increase the diversity of SSEs.

Result and Discussion
As shown in Figure 1a, Li 15 P 4 S 16 Cl 3 adopts the space group of I43 d with two Li sites (12a and 48e).Those Li sites are coresharing, and the migration of Li þ ions between adjacent sites will pass through the intermediate octahedral sites (Figure 1b).The Li þ ion migration barrier along this path (T-O-T in Figure 1b) is larger than 0.4 eV due to the high site energy around octahedral site (Li3, 48e). [11]Since Li þ ions are more stable at octahedral location than tetrahedral sites for halide, [36] being opposite to sulfides, the mixing of halide/S anion is expected to effectively reduce the energy difference between tetrahedral sites and octahedral sites.However, the high stability of PS 4 makes it difficult to dope halides at the S site.As shown in Figure S1, Supporting Information, even a small amount of Cl À doping (Li 14.8 P 4 S 15.8 Cl 3.2 ) can bring unknown impurities.Instead, O 2À was herein selected as the dopant to regulate the energy difference between tetrahedral and octahedral sites in Li 15 P 4 S 16 Cl 3 .Previous results indicate that the doping of O 2À in sulfides helps regulate site energy and reduce the Li þ ion migration activation energy. [30,31]xygen-doped Li 15 P 4 S 16Àx O x Cl 3 (x = 0, 0.2, 0.4, 0.6) was synthesized via solid-state reaction.Unlike halogen doping, the introduction of O 2À ions at the S site is more stable.No obvious impurity is observed in the XRD even for the sample with O 2À content up to 0.6 (Li 15 P 4 S 15.4 O 0.6 Cl 3 ; Figure 2a), indicating successful doping of O 2À .Furthermore, Raman spectra were collected and shown in Figure S2, Supporting Information.The peak at around 419 cm À1 can be assigned to PS 4 3À and the peak at around 388 cm À1 should be the O 2À -doped PS 3 O 3À , since it is located between 419 cm À1 (PS 4 3À ) [37] and 360 cm À1 (PO 4 3À ). [38]The area ratio between the peak at 388 and 419 cm À1 was analyzed, which keeps increasing with higher O 2À content (1:12, 1:7.6, and 1:6.4 with 0.2, 0.4, and 0.6 O 2À , respectively).Those results also confirm the doping of O 2À .Electrochemical impedance spectroscopy (EIS) in Figure 2b reveals a low ionic conductivity of 1.6 Â 10 À9 S cm À1 for the undoped Li 15 P 4 S 16 Cl 3 , being sorted as a nonionic conductor.The ionic conductivity can be enhanced to the highest value Although an appropriate Li þ /O 2À ion radius ratio allows Li þ ions to be possibly stable at both tetrahedral and octahedral sites of oxides, limited O 2À in Li 15 P 4 S 16Àx O x Cl 3 makes the octahedral site still unstable for Li þ , thus the activation energy is only slightly reduced.As a result, the ionic conductivity enhancement is limited, which still cannot meet the requirements for SSEs (Figure 2d).
To further increase the ionic conductivity in Li 15 P 4 S 16 Cl 3 , the sintered sample was further ball milled to introduce local structure distortion, such as a change of bond length or bond angle, which would change the bond energy and thus Li þ site energy.The ball milling speed was fixed at 500 rpm and the duration time was varied from 15 to 60 min (Figure S4 and S5, Supporting Information).Peak intensities in XRD patterns gradually decrease with the increase of ball milling time, indicating a decrease in the crystallinity.Transmission electron microscopy (TEM) images in Figure S6, Supporting Information, show that the sample maintains mostly crystalline and only a small amount transits to the amorphous phase after 30 min ball milling, indicating that the decrease in crystallinity is mainly due to changes in short-range structures, such as distortion.The ionic conductivity is significantly enhanced to 5.58 Â 10 À5 S cm À1 and activation energy is reduced to 0.33 eV with optimized ball milling time (30 min).When longer ball milling time is applied and more amorphous phase is introduced, the ionic conductivities decrease to 4.2 Â 10 À5 and even to 2.2 Â 10 À5 S cm À1 with 45 and 60 min ball milling.This indicates that the amorphous phase here is not the main conducting phase and a certain ordered structure is needed to achieve high ion conductivity.
Since the local structure distortion has a more significant effect on improving ionic conductivity, ball milling (30 min) is combined with O 2À doping to further enhance the ionic conductivity of Li   3c), according to the variable temperature EIS (Figure S7, Supporting Information).Combining oxygen doping and local structure distortion, the ionic conductivity of Li 15 P 4 S 16 Cl 3 can be enhanced from 10 À9 to 10 À4 S cm À1 level, being served as a fast ion conductor.Anion mixing (doping) presents wide access to most inorganic SSEs for enhancing conductivity.However, ball milling strategy used here shows potential universality for soft inorganic SSEs, such as sulfides and halides.In addition to ionic conductivity, the electronic conductivity of the ball-milled Li 15 P 4 S 15.6 O 0.4 Cl 3 is measured as 1.55 Â 10 À10 S cm À1 , indicating a pure ion conductor (Figure S8, Supporting Information).
According to the previous analysis, O 2À doping and ball milling can effectively regulate the Li þ site energies, thereby achieving lower ion migration activation energy.Meanwhile, changes in site energy will alter the distribution of Li þ ions in Li 15 P 4 S 16Àx O x Cl 3 .As shown in Figure 4a, solid-state 6 1b), the more S ion coordination leads to the left shift of Li peak. [39]The distortion in Li 15 P 4 S 16 Cl 3 after ball milling changes the site energy and stability of Li þ ions, resulting in the migration of Li þ ions from LiS 4 to LiS 6 site.When combining O 2À doping and ball milling, five different Li coordinations all show up in the 6 Li NMR spectra.
Tracer-exchange NMR is a powerful method for detecting ion migration pathways in solid electrolytes. [5,40,41]To explore the mechanism of O 2À doping and local structure distortion on the increase of ionic conductivity, symmetric 6 Li| Li 15 P 4 S 16Àx O x Cl 3 | 6 Li cells were assembled and polarized.The 6 Li atoms will replace the 7 Li and/or 6 Li on their way upon polarization, which will make the difference in 6 Li NMR spectra, as shown in Figure 4b.A significant increase in the intensity of LiS 3 Cl and LiS 5 Cl after polarization, with the signal strength of LiS 4 did not change significantly, revealing that the main  5).For pristine Li 15 P 4 S 16 Cl 3 , the electrochemical window is measured to be from 2.10 to 2.69 V (versus Li-Si).Compared to typical sulfide SSEs (≈1 V), [42][43][44] the reduction potential is much higher, making this SSE not ideal when paired with low-potential anodes, such as Li metal.After O doping, the oxidation stability  The electrochemical window given by cyclic voltammetry is based on thermodynamic stability, which is quite narrow for sulfide-based SSEs.However, in practical applications, the low electronic conductivity of the oxidation or reduction products of sulfide solid electrolytes can stabilize the interface between them, allowing for dynamic stability over a wider voltage range. [1,9,45,46]When cycling at low current and low capacity density, sulfide SSEs stabilize in Li|SSE|Li symmetric cells, with no increase in working voltage and interfacial resistance. [9]At high currents, the instability of Li|SSE interface is mainly due to the formation of voids, which reduces the physical contact between Li and SSE. [9]Here, the dynamic stabilities of Li 15 P 4 S 16 Cl 3 and Li 15 P 4 S 15.6 O 0.4 Cl 3 against Li metal were also evaluated in Li|SSE| Li symmetric cells, with a small current density of 0.01 mA cm À2 and capacity density of 0.002 mAh cm À2 (Figure 6a).Before cycling, the cell of Li|Li The all solid state batteries (ASSBs) using Li 15 P 4 S 16 Cl 3 and Li 15 P 4 S 15.6 O 0.4 Cl 3 as SSE, Li-Si as anode, and LiCoO 2 as cathode were assembled and cycled at 25 °C (Figure 6b-d).Here, siliconbased anode (Li-Si) was selected because of its higher potential [47] and better stability to SSE (Figure 6b).Before cycling, EIS for different batteries was tested to confirm whether the battery works well.As shown in Figure S10

Conclusion
In summary, a combined strategy with anion mixing (oxygen) and local structure distortion (ball milling) is applied to Li 15 P 4 S 16 Cl 3 to improve the ionic conductivity.In ordered Li 15 P 4 S 16 Cl 3 , Li þ ions are distributed in two core-sharing tetragonal sites (12a, 48e).Lithium ions migrate through the path of tetragonal (stable), interstitial octahedral (unstable), and then o tetragonal (stable) in Li 15 P 4 S 16 Cl 3 , with a high activation energy of 0.50 eV and low ionic conductivity of 1.6 Â 10 À9 S cm À1 .With O 2À doping and local structure distortion, changes in local

Experimental Section
Material Synthesis: Li 15 P 4 S 16Àx O x Cl 3 (x = 0, 0.2, 0.4, 0.6) was synthesized via solid-state reaction, using Li 2 S (99%, Macklin), P 2 S 5 (99%, Macklin), LiCl (99%, Macklin), and P 2 O 5 (99.9%,Aladdin) as raw materials.The raw materials were weighed in a stoichiometric ratio and manually grounded first.Then the mixture was sealed in a zirconia jar and ball milled at 500 rpm for 300 min on a planetary ball milling machine (Nanda Instrument, QM-3SP2).After that, the obtained Li 15 P 4 S 16Àx O x Cl 3 precursor powder was sealed in quartz tubes and sintered at 400°C for 20 h, with a heating rate of 5°C min À1 .Finally, the sintered samples were sealed and ball-milled at 500 rpm again to obtain the final products.
Material Characterizations: The phase purity and crystal structure Li 15 P 4 S 16Àx O x Cl 3 electrolytes were characterized by X-ray diffraction (XRD, PANalytical X-Pert PRO MPD) using Cu Kα radiation (λ = 1.5406Å) with a 2θ range of 10°-80°.To avoid air exposure, samples were covered by Kapton film when collecting XRD data.Raman spectra were collected with a microconfocal laser Raman spectrometer (HORIBA JOBIN YVON), which was covered with glass to avoid exposure to air. 6Li solid-state NMR experiments were performed on a Bruker 600M spectrometer (14.1T) with AVANCCE NEO consoles using a Bruker 3.2 mm HXY MAS probe.The samples were filled into rotors inside the Argon glove box.The Larmor frequencies for 6 Li were 88.32 MHz, respectively. 6Li spectra were acquired by using one-pulse program with a small flip angle (<π/6) and were referenced to 1 M LiCl solutions with chemical shifts at 0 ppm.The spinning rate ν rot was set to 14 kHz.TEM images were acquired using JEM-2100 F.
Electrochemical Measurements: The ionic conductivities and activation energies of the electrolyte materials were determined by variable temperature EIS with an electrochemical workstation (Bio-Logic SP-200) in the frequency range of 5 MHz-1 Hz.The powders were pressed into pellets with a diameter of 12 mm and thickness of around 1 mm, at a pressure of 160 MPa.
(7.0 Â 10 À9 S cm À1 ) with oxygen doping in Li 15 P 4 S 15.6 O 0.4 Cl 3 .Further increasing the proportion of O 2À could introduce some unknown impurities, which reduces the ionic conductivity to 4.0 Â 10 À9 S cm À1 (Li 15 P 4 S 15.4 O 0.6 Cl 3 ).Activation energies were determined from variable temperature EIS measurements (Figure S3, Supporting Information) with 0.50 eV for the undoped Li 15 P 4 S 16 Cl 3 (Figure 2c).After oxygen doping, the activation energy decreases to 0.43 eV for Li 15 P 4 S 15.8 O 0.2 Cl 3 , 0.47 eV for Li 15 P 4 S 15.6 O 0.4 Cl 3 , and 0.45 eV for Li 15 P 4 S 15.4 O 0.6 Cl 3 , respectively.
15 P 4 S 16 Cl 3 .As shown in Figure 3a, the intensities are significantly reduced but still can be indexed to pure cubic phase.The Nyquist plots of ball-milled Li 15 P 4 S 16Àx O x Cl 3 (x = 0, 0.2, 0.4, 0.6) at room temperature are shown in Figure 3b, in which Li 15 P 4 S 15.6 O 0.4 Cl 3 shows the lowest impedance, corresponding to the highest ionic conductivity of 1.15 Â 10 À4 S cm À1 .Meanwhile, activation energies are reduced to 0.32 eV (Li 15 P 4 S 16 Cl 3 ), 0.34 eV (Li 15 P 4 S 15.8 O 0.2 Cl 3 ), 0.35 eV

Figure 1 .
Figure 1.a) Crystal structure and b) typical Li þ ion migration path in Li 15 P 4 S 16 Cl 3 .
Li magic angle spinning (MAS) nuclear magnetic resonance (NMR) spectra reveal two Li þ ions at 48e (1.32 ppm, LiS 3 Cl) and Li þ at 12a (À0.55 ppm, LiS 4 ) in Li 15 P 4 S 16 Cl 3 .After doping with 0.4 O 2À , extra peaks at around 1.66 ppm and À1.18 ppm show up, corresponding to LiOS 2 Cl and LiCl, respectively.The formation of LiOS 2 Cl is caused by the substitution of O for S in LiS 3 Cl.Since the electronegativity of O is larger than S, the resonance of LiOS 2 Cl appears at the low field.There is no extra peak presents at the left of LiS 4 , revealing that S (48e) around LiS 4 is hardly replaced by O and the S at 16c site is more easily replaced by O.After 30 min ball milling, a new signal of Li is generated at 0.6 ppm, accompanied by the disappearance of the signal of LiS 4 .This new signal at 0.6 ppm is from Li at LiS 5 Cl octahedral site (Li3 shown in Figure

Figure 2 .
Figure 2. a) Powder X-ray diffraction (PXRD) patterns of Li 15 P 4 S 16Àx O x Cl 3 (x = 0, 0.2, 0.4, 0.6) after high-temperature sintering.b) The Nyquist plots of Li 15 P 4 S 16Àx O x Cl 3 were measured at room temperature.c) The Arrhenius plots of Li 15 P 4 S 16Àx O x Cl 3 .d) Ionic conductivities and activation energies of Li 15 P 4 S 16Àx O x Cl 3 .

Figure 3 .
Figure 3. a) PXRD patterns of Li 15 P 4 S 16Àx O x Cl 3 (x = 0, 0.2, 0.4, 0.6) after secondary ball milling.b) The Nyquist plots of Li 15 P 4 S 16Àx O x Cl 3 .c) The Arrhenius plots of Li 15 P 4 S 16Àx O x Cl 3 .d) Ionic conductivities and activation energies of Li 15 P 4 S 16Àx O x Cl 3 after secondary ball milling.

Figure 4 .
Figure 4. a) Solid-state6 Li MAS NMR spectra of Li 15 P 4 S 16Àx O x Cl 3 samples and b) the spectral comparison and difference before and after polarization with metallic6 Li electrodes.
15 P 4 S 16 Cl 3 |Li gives a total resistance over 5000 Ω, in which Li 15 P 4 S 16 Cl 3 SSE contributes 3800 Ω and Li 15 P 4 S 16 Cl 3 /Li metal interface contributes 1500 Ω.During cycling, the voltage slightly increases from 0.01 V (at the beginning) to around 0.015 V (22 h).After that, the Li|Li 15 P 4 S 16 Cl 3 |Li symmetric cell is short-circuited, indicating the generating of dendritic lithium (Figure 6a, black line).For the cell of Li|Li 15 P 4 S 15.6 O 0.4 Cl 3 |Li, Li 15 P 4 S 15.6 O 0.4 Cl 3 SSE contributes 800 Ω and Li 15 P 4 S 15.6 O 0.4 Cl 3 /Li metal interface contributes 350 Ω resistance before cycling (Figure S9, Supporting Information).The voltage of Li|Li 15 P 4 S 15.6 O 0.4 Cl 3 |Li at the beginning is a little bit lower than 0.01 V and increases to around 0.015 V after 1000 h cycles, indicating high stability of Li 15 P 4 S 15.6 O 0.4 Cl 3 against Li metal.
These pellets were sandwiched by two stainless steel rods and EIS data were collected from 25 to 100 °C.To evaluate the electrochemical stability, cyclic voltammetry measurements were carried out on Li-Si|Li 15 P4S 16Àx O x Cl 3 |SSE þ C between open circuit voltage and 0 V, or between open circuit voltage and 5 V at a scanning rate of 0.5 mV s À1 .Battery Assembly and Characterization: The all-solid-state batteries were assembled with LiCoO 2 composite cathode, Li 15 P 4 S 16 Cl 3 or Li 15 P 4 S 15.6 O 0.4 Cl 3 SSEs, and Li-Si anode.The LiCoO 2 composite cathode is composed of LiCoO 2 , SSE, and Super P in a weight ratio of 5:5:0.1.The mixed cathode was hand-ground in a mortar and pestle for 10 min and then placed in a stainless-steel jar and ball milled at 200 rpm for 1 h.Li-Si powder is a mixture of silicon powder and fine lithium flakes in a molar ratio of 1:1, obtained by ball milling in a stainless-steel ball milling jar at 550 rpm for 10 h.30 mg Li-Si anode, 50 mg SSE, and 8 mg LiCoO 2 composite cathode were placed and prepressed under 80 MPa in a homemade battery mold one by one.Finally, the battery was pressed under a pressure of 160 MPa to obtain pellet-type ASSB with a diameter of 10 mm.The ASSBs were cycled at room temperature and 0.5 C (1 C = 140 mA g À1 ) within the voltage range of 2.5-4.0V, using a NEWARE battery testing system.
65 P 4 S 16 Cl 3 as a model material to adjust the site energy of Li þ ions at different sites (tetrahedral and DOI: 10.1002/sstr.202300565Highionicconductivity is the key point in the development of new solid-state electrolytes.Herein, a combining strategy of anion (O 2À ) doping and structure distortion is applied to enhance the Li þ ion conductivity in Li 15 P 4 S 16 Cl 3 , thus converting the nonionic conductor into fast ionic conductor.Solid-state6Li nuclear magnetic resonance analysis shows redistribution of Li þ ions in Li 15 P 4 S 16 Cl 3 with O 2À doping or local structure distortion via ball milling, indicating energy changes at different lithium sites.As a result, the activation energy is reduced from 0.50 to 0.35 eV for the ball-milled Li 15 P 4 S 15.6 O 0.4 Cl 3 , and the ionic conductivity is enhanced from 10 À9 to 10 À4 S cm À1 .The electrochemical stability of Li 15 P 4 S 15.6 O 0.4 Cl 3 is broadened at the anode side as well.The symmetric cell Li| Li 15 P 4 S 15.6 O 0.4 Cl 3 |Li can cycle more than 1000 h with negligible voltage increase.The LiCoO 2 |Li 15 P 4 S 15.6 O 0.4 Cl 3 |Li-Si all-solid-state battery demonstrates an initial capacity of 106 mA h g À1 and retains 92% capacity after 200 cycles at 0.5 C, highlighting excellent rate performance and electrochemical stability.