Dr L Pasquini

R.-A Becerra-Arciniegas

Dr R Narducci

Dr E Sgreccia

V Gressel

Prof. M L Di Vona

Prof P Knauth

Properties and alkaline stability of composite anion conducting ionomers based on poly(phenylene oxide) grafted with DABCO and Mg/Al lamellar double hydroxide

Keywords: degradation, ion exchange capacity, ionic conductivity, membranes, nanoparticles

We studied composite anion exchange membranes made from poly(2,6-dimethyl-1,4phenylene oxide) (PPO) with grafted 1,4-diazabicyclo[2.2.2]octane (DABCO) and Mg/Al lamellar double hydroxide particles as inorganic filler. The mechanical and conductivity properties were measured in fully humidified conditions. The activation energy for conductivity amounts to (0.20 ± 0.01) eV. The degradation of the anion exchange membrane in 2 M NaOH at 60 °C was investigated during 168 h following the ionic conductivity, water uptake and ion exchange capacity as function of the alkaline treatment time. The degradation is rapid within the first hours and then slows down. The degradation was also studied by thermogravimetric analysis and FTIR spectroscopy, showing the various decomposition phenomena of the composite.

Introduction

The mediocre alkaline stability is the main obstacle contrasting the widespread application of anion exchange membranes in alkaline fuel cells and other electrochemical devices [1][2][3][4][5][6][7][8].

However, the availability of such membranes would allow avoiding the extensive use of platinum or related expensive noble-metal electrocatalysts for the oxygen reduction reaction [9][10][11]. The alkaline degradation can occur on functional groups, especially quaternary ammonium moieties [12][13][14][15], but also on the polymer backbone [3,16], especially if it contains electron-withdrawing groups, such as sulfone. Polysulfones (PSU) [17,[START_REF] El-Hibri | Polysulfones[END_REF] are therefore considered less alkaline stable polymers and ionomers based on poly(2,6-dimethyl-1,4-phenylene oxide) (PPO), which do not contain sulfone links, are expected to present a higher alkaline stability [15,[START_REF] Pan | [END_REF][20][21][22][23][24], although the ether groups are still possible weak points [25][26][27].

We have previously reported the beneficial effect of bulky and stiff amines, such as 1,4diazabicyclo[2.2.2]octane (DABCO), which can mitigate the degradation effect by steric hindrance [5]. In the present work, we use quaternary ammonium groups made with DABCO grafted on PPO chains. The ionomer itself should therefore present a decent alkaline stability. Furthermore, we want to explore the synergistic effect of the formation of composites, especially using an inorganic solid that contains itself mobile anions, which can partly compensate the loss of ionic conductivity due to the presence of an inorganic filler [28,29].

As recently reported, inorganic layered compounds in mixed matrix membranes improve the performance of AEM [30,31]. Lamellar double hydroxides (LDH) [32] are layered ionic conducting inorganic solids that contain mobile anions in the interlamellar space, which can be exchanged and can help to maintain a certain level of conductivity even in low humidity environment [33]. LDH show exceptionally high and anisotropic 2D ion conduction [34,35] via the presence in the layered structure of linked oxygenated groups allowing the establishment of a hydrogen bond network along the LDH surface. They might also mitigate the attack of anion exchange membranes in alkaline conditions by composite effects and reduce the membrane swelling and excessive water uptake. LDH can be prepared with various doping cations. We have previously reported the growth of Zn-Al LDHs in thin film form [36]. We have also studied Mg-Al [37] and Zn/Al [38] LDH as second phase in composites of PSU with grafted quaternary ammonium groups.

Consequently, we study in this work the properties and the alkaline stability of composite PPO-based ionomers containing quaternary ammonium groups with DABCO and Mg-Al LDH as inorganic phase. The LDH composition was Mg 0.66 Al 0.33 (OH) 2 (NO 3 ) 0.33 . The mobile nitrate ions can be exchanged against chloride or hydroxide ions.

The concentration of the inorganic compound is quite high (30 wt%) to explore the beneficial effects on the membrane water uptake and degradation. The purpose of a very high loading is threefold: i) get as close as possible to a mainly inorganic membrane sustaining good mechanical properties, ii) minimize the degradation due to the loss of ionomeric groups, and iii) preserve the ionic conductivity of the composite membrane by creation of a percolating anion conducting inorganic part.

Experimental

Products

Poly(2,6-dimethyl-1,4-phenylene oxide) (PPO, MW = 50000 g/mol) and all reagents, such as 1,4-diazabicyclo[2.2.2]octane (DABCO), paraformaldehyde, trimethylchlorosilane, SnCl 4 , Mg(NO 3 ) 2 .6H 2 O and Al(NO 3 ) 3 .9H 2 O, were reagent-grade and purchased from Sigma-Aldrich.

Ionomer synthesis

The ionomer was prepared as reported previously [24]. In short, PPO was first chloromethylated and subsequently quaternized by nucleophilic substitution with DABCO [38]. PPO (5.04 g, 42 mmol) was dissolved in 500 mL chloroform and 198 mmol (5.98 g) of paraformaldehyde was introduced and stirred during 1 h. Then 4.2 mmol (0.46 mL) of SnCl 4 was added at room temperature. A solution of trimethylchlorosilane (198 mmol, 25 mL) in chloroform was introduced dropwise during 2 h at RT. To obtain a degree of chloromethylation DCl = 0.28, the reaction temperature was kept at 45 °C for 17 h. After this time, the temperature was decreased and 150 mL of methanol added at RT until complete precipitation. The formed Cl-PPO was washed with methanol and dried in an oven at 60 °C for 48 h.

The chloromethylated precursor was then quaternized in 1-methyl-2-pyrrolidinone (NMP) by reaction with an excess of DABCO to avoid the possible cross-linking reaction due to the presence of two nitrogens in the structure [1]. Typical quantities of reactants were 12.7 mmol (1.68 g) of Cl-PPO (DCl = 0.28) and 4.9 mmol of DABCO (0.56 g) in 125 mL NMP, corresponding to a molar excess of 1.4 with respect to the functional ammonium groups; this means an excess of 10 mol% with respect to the polymer. The reaction was carried out during 72 h at 70 °C under nitrogen flux. After this time, the excess of DABCO was eliminated by heating the solution for 2 h at 80 °C under vacuum using a liquid nitrogen trap and a rotary pump (2 10 -6 bar). The remaining solution (about one third of the initial volume) was directly used for the casting procedure. The prepared membranes were repeatedly washed in deionized water to remove any remaining traces of DABCO, which is water-soluble. The membranes were obtained in Cl form. The degree of functionalization (DF = 0.27), measured by Mohr titration (see below), was nearly identical to the initial DCl.

The structure of PPO-DABCO polymer in Cl form is shown in Figure 1.

OH O O O H N + N n Cl - Figure 1. PPO-DABCO structure.

LDH synthesis

Mg/Al-LDH in nitrate form was synthesized by a co-precipitation method with a molar ratio Mg/Al = 2 corresponding to a formula Mg 0.66 Al 0.33 (OH) 2 (NO 3 ) 0.33 . De-carbonated water was used to prepare the two salt solutions and avoid the formation of Mg/Al carbonates. 0.02 mol Mg(NO 3 ) 2 .6H 2 O and 0.01 mol Al(NO 3 ) 3 .9H 2 O were slowly added in a flask at room temperature under magnetic stirring under a constant flow of N 2 gas. A solution of 1 M NaOH was then added dropwise under magnetic stirring until pH = 10.0 was attained. The pH was controlled by in situ potentiometric titration. The resulting slurry was heated for 18 h at 80 °C.

The white precipitate was washed with deionized water many times and dried in an oven at 80 °C for 48 h.

Preparation of composite membranes

PPO-DABCO composite membranes with 30 wt% LDH were prepared by adding the filler to 20 mL of a 0.05 M solution of PPO-DABCO in NMP. The resulting mixture was stirred for 4 h and evaporated to around 10 mL on a heating plate at 100 °C. The viscous solution was then cast onto a Petri dish and dried in an oven at 100 °C during 48 h. The composite membranes contain initially Cl -ions from the synthesis of PPO-DABCO, and NO 3 -ions from the synthesis of LDH. For this reason, the counterions were systematically exchanged with Cl -or OH -ions before any measurements (see below).

Ion Exchange Capacity (IEC) by Mohr titration

This method is based on ion exchange reactions and determines the chloride ion concentration In this work, a 0.01 M AgNO 3 solution was prepared for the Mohr titration avoiding sunlight.

The IEC is expressed as the quantity of anions that can be exchanged by a specific mass of dry ionomer:

(3)

Water uptake

The water uptake WU was measured for both Cl and OH forms. The membranes in Cl form were obtained by immersion in 1 M NaCl solution during 48 h then carefully washed in deionized water. The OH form was obtained by immersion of the membranes in 2 M KOH during 48 h to exchange Cl -ions and washing them 3 times in deionized water. Manipulations were done ensuring that the membranes were as little as possible in contact with air avoiding carbonatation problems.

The water uptake is defined by the mass of wet sample (m wet ) and the mass of dry sample

(m dry ): (4)
To determine the wet mass, samples were removed from the deionized water and carefully wiped with adsorbing paper. To determine the dry mass, the samples were stored in a desiccator over P 2 O 5 during 48 h to remove completely the water.

The Water Uptake WU (in %) and IEC (in meq/g) permit to calculate the hydration number λ, which represents the number of water molecules per quaternary ammonium group:

Mechanical measurements

Tensile stress-strain tests were performed using an Adamel-Lhomargy DY30 uniaxial traction machine at 25 °C on fully humidified membranes in Cl form with a traction rate of 5 mm/min.

After immersion of the membranes in deionized water for 24 h, in order to reach the maximum humidification, samples of 5 mm width and 25 mm length were first carefully cut in the centre of the membranes to remove inhomogeneous parts at the borders and then rapidly installed between the clamps of the machine. The short duration of the test allows maintaining full humidification during the experiment.

Thermogravimetry

The thermogravimetric experiments were performed using a TA Instruments TGA Q500 high resolution analyser. Around 7 mg of Cl form ionomers were placed in Pt sample holders and heated with a maximum rate of 3 K/min from 30 to 600 °C in air. Comparative thermograms were recorded after degradation of the ionomers in 2 M NaOH at 60 °C for 168 h.

FTIR spectroscopy

FTIR spectra were recorded between 4000 and 450 cm -1 using a Perkin Elmer Spectrum 2 IR spectrometer equipped with an ATR crystal diamond module. The membrane was directly analysed by squeezing it on the diamond window at room temperature and humidity in Cl form.

Ionic conductivity by impedance spectroscopy

The sample resistance impedance of Cl form membranes was measured at full humidification.

The samples in Cl form were first immersed in deionized water during 24 h; after removing excess water on the surface with absorbing paper, they were immediately introduced inside a hermetically closed Swagelok cell with stainless steel electrodes. The analysis was performed between 1 Hz and 6 MHz with a signal amplitude of 20 mV using a VSP-300 Electrochemical Impedance Spectrometer (Biologic science instruments). The resistance of the membranes in

Results and discussion

Mechanical properties

Figure 2 shows a typical mechanical stress-strain test. The mechanical properties were determined in fully humidified environment to simulate the operative condition of membranes. The Young modulus (E = 280 ± 20 MPa) and tensile strength (TS = 9 ± 2 MPa) are consistent with a highly hydrated membrane, because the high dielectric constant of water reduces the Van der Waals interactions between ionomer chains [39,40]. The formation of composites leads to a slight increase of the Young modulus (E = 320 ± 60 MPa); the tensile strength decreases instead (TS = 3 ± 2 MPa). The elongation at break is quite low (< 5 %) in all cases.

Temperature dependence of the ionic conductivity

The chloride ion conductivities of PPO-DABCO and PPO-DABCO + LDH between 25 and 80 °C are reported in Table 1 and plotted vs the reciprocal absolute temperature in Figure 3.

PPO-DABCO PPO-DABCO + LDH

The presence of the inorganic phase reduces the ionic conductivity, as previously observed in other composites [37], but the decrease is relatively limited also considering the high quantity of the filler in the membranes. The low decrease observed is attributable to the presence in the LDH structure of mobile anions that contribute to the conductivity of samples, although they are not enough to compensate the loss of conductivity. One must also take into account the necessary interconnection between two-dimensional ion-conduction planes in LDH and onedimensional ion-conduction channels in the ionomer limiting the percolative anion transport through the membranes. Below the glass transition temperature of the ionomer (which is nearly 200 °C for related PSU-DABCO [39]), the activation energy Ea for chloride ion conduction can be calculated using an Arrhenius-type equation: (7) σ 0 is the prefactor, R is Boltzmann's constant and T is the absolute temperature. The activation energy, calculated from the graph in Figure 3, is very similar for both materials and amounts to (0.20 ± 0.01) eV in good agreement with other ionomers in Cl form [41]. Figure 4 shows the ionic conductivity of pristine and composite membranes after the aging in KOH at different times. The ionic conductivity decrease of PPO-DABCO is similar to previously reported data for short side-chain ionomers (23% of residual conductivity after 168

Degradation test: ionic conductivity

PPO-DABCO PPO-DABCO + LDH

h of alkaline treatment) [21]. The composite membrane shows a comparable or slightly better degradation resistance. One should mention that the alkaline degradation is strongly correlated with the hydration. For low hydration numbers, typically below 4, the incomplete solvation shell of the hydroxide ions increases their reactivity and amplifies the alkaline degradation [42,43].

Degradation test: ion exchange capacity

A similar behaviour can be noticed observing the ion exchange capacity dependence on the time of aging obtained from Mohr titration (Figure 5). The IEC decrease is correlated with the loss of ammonium groups in alkaline conditions. The strongest degradation is observed within the first 12 h of the treatment. At longer times, the degradation is much slower. We have previously attributed this result to a pseudo second-order kinetics of the degradation reaction [24].

Figure 5. Ion exchange capacity as function of the alkaline treatment time.

Degradation test: water uptake

Water uptake data for OH and Cl forms are shown in Figure 6. One can notice the about twice higher water uptake for the OH form in agreement with the expectation, given that hydroxide anions are much more hydrated than chloride anions [44]. The decrease with alkaline treatment time is relatively low.

In the composite, the water uptake is slightly lower for both OH and Cl forms, although the difference is not large. Especially in the OH form, WU increases with the time of the aging, apparently in contradiction with the decrease of the IEC. A possible explanation is that the aging in alkaline condition of the composite membranes for a long time acts on LDH as a post synthetic hydrothermal treatment. It is reported in the literature that the hydrothermal aging may lead to particles with larger size altering the absorption of water [45,46]. This effect might justify the observed larger scatter for the composite. The evolution of the hydration number  for the composite membrane is reported in Figure 7 and indicates that the change in the LDH morphology starts after 60 h in KOH. observed at a similar temperature [21]. The corresponding mass loss of about 15 % is slightly below the calculated value. A second sharp mass loss of 18 % around 360 °C is attributed to the loss of the two methyl groups of PPO, which decompose prior to the main chain [47]; the calculated mass loss is in excellent agreement with the experiment. The remaining mass above 380 °C (62 %) is in very good agreement with the calculated value for the remaining polymer backbone. For degraded PPO-DABCO+LDH, the mass loss above 130 °C corresponds to the elimination of interlayer structural water [48,49]; the following decomposition above 200 °C is very diffuse with a continuously sloping curve. It is impossible to distinguish the various effects and the differential thermogravimetric curves (DTG) are more useful in this respect (see below). The residual mass of around 16 % at the end of the experiment belongs to oxides due to decomposition of the LDH; the difference with the initial percentage of LDH (30 %) is due to the loss of anions, as previously described [48, To improve the analysis, the differential thermogravimetric curves (DTG) are more useful:

each peak corresponds to a thermal degradation process. For PPO-DABCO (Figure 8b) after the initial mass loss due to membrane dehydration up to 100 °C, the first peak with a maximum at 230°C is related to the loss of DABCO. One can notice that the peak is lower after the alkaline treatment, where a part of the quaternary ammonium moieties is lost. The second peak with a maximum near 360°C, attributed to the methyl groups of the PPO backbone [24], is very significantly reduced after the alkaline treatment. Ehlers reported that the decomposition mechanism of PPO begins with the cleavage of the ether linkages and simultaneous removal of the methyl groups [47]; probably a similar mechanism is acting during the alkaline degradation. The final large peak with a maximum near 415 °C corresponds to the decomposition of the remaining polymer backbone.

In PPO-DABCO+LDH composites (Figure 8c), the first peak with a maximum near 150 °C belongs to removal of interlayer water of LDH [48,50]. In the broad signal starting around 200 °C, one can first observed the loss of DABCO with an intermediate maximum at 230 °C.

At higher temperature, the decomposition reactions of LDH overlap the decomposition of the polymer. They include the loss of intercalated structural water, dehydroxylation of the brucite-like layer and removal of interlayer chloride ions, giving a mixture of oxides [50].

These processes in the LDH structure were reported at comparable temperatures in the literature, e. g. above 280 °C and up to 500 °C with a total weight loss of 45% [50] and with a maximum at 310 [49] or 320 °C [48]. One can notice that the alkaline treatment strongly modified the aspect of the DTG curve. The smaller and broader peaks are related to more continuous decomposition processes (cf. Figure 8a). Evidently, LDH is modified during the alkaline treatment and some reactions may occur, possibly related to the amphoteric nature of aluminium oxide. The peak of the PPO main chain degradation has a maximum at 410 °C, slightly below PPO-DABCO, probably due to the interaction with the inorganic component.

After the alkaline treatment, a second peak with a maximum at 420 °C is observed, corresponding to further decomposition processes of LDH, which continue up to 500 °C as previously demonstrated [50]. The intensity of the OH stretching absorption around 3400 cm -1 is lower for the aged sample, which may be related to the loss of hydrophilic ammonium groups after the KOH treatment.

FTIR spectroscopic analysis

Conclusions

The idea of this work was to explore the effect of inorganic phase particles in large amount dispersed in an ionomer matrix in order to improve the alkaline and hydrolytic stability of anion exchange membranes. The second phase was a Mg/Al lamellar double hydroxide, which was itself anion conducting. AEM could be prepared successfully and their conductivity was relatively slightly reduced in comparison with the pristine ionomer. The degradation behaviour, studied by ex situ conductivity, water uptake and ion exchange capacity measurements, was also relatively similar to the pristine ionomer with a relatively fast degradation during the first hours. Thermogravimetric experiments and FTIR spectra made before and after the alkaline degradation gave complementary information on the degradation processes.

Evidently, the approach of a high anion-conducting filler content has not improved sufficiently the alkaline resistance of the membranes and the other key properties are also not significantly enhanced. Future work might instead be focused on a better dispersion of the inorganic compound into the ionomer, for example by in situ sol-gel synthesis or similar means, and by a more clever exploitation of the 2D morphology of LDH particles.

Table of Contents Graphic

Composite anion exchange membranes made of a chemically resistant polymer backbone poly(phenylene oxide) (PPO), a stiff and bulky functionalizing amine (DABCO) and an anion conducting layered double hydroxide (LDH) as filler are a promising combination to reduce the instability in alkaline solution. The degradation kinetics is followed by measurements of the ionic conductivity, the ion exchange capacity and the water uptake. Its mechanism is studied by thermogravimetric analysis and FTIR spectroscopy.

2 -; 1

 21 by titration with silver nitrate. The anions, initially present in PPO-LDH samples, were exchanged by Cl -in a 1 M NaCl solution during 48 h at room temperature. Then, membranes were washed in deionized water for 30 minutes and transferred into 20 mL of 1 M Na 2 SO 4 solution to exchange Cl -by SO 4 mL of 0.25 M K 2 CrO 4 indicator solution was added.As the AgNO 3 solution was slowly introduced, a precipitate of AgCl formed (equation 1). The end-point of the titration was reached when all the chloride ions were precipitated. Then additional silver ions reacted with the indicator, potassium chromate, to form a red-brown precipitate (equation 2).

6 ) 2 . 11 .

 6211 fully humidified form was obtained at 25, 40, 60 and 80 °C from a Nyquist plot by non-linear least-square fitting. The equivalent circuit was a classical series arrangement of two resistance-constant phase elements in parallel. The ionic conductivity σ was calculated using equation (6) with the resistance R, the membrane thickness d, typically 60-80 μm, and the electrode area A = 0.264 cm 2 . σ = d / (A*R) (Alkaline stability tests To carry out the stability tests, 8 samples of each membrane were cut in rectangles of approximately 40 mg mass. First, the membranes were washed by immersing them in deionized water for 48 h in an oven at 60 °C. The stability tests were performed by putting the samples into 2 M KOH at 60°C for various times between 2 and 168 h inside a Teflon container fluxed with nitrogen. The samples were then transformed into the Cl form by immersion in 2 M NaCl solution during 48 h and then washed with deionized water. The sample conductivity, IEC and WU were then determined as described above.

Figure 2 .

 2 Figure 2. Tensile stress-strain test for PPO-DABCO in Cl form at full humidification at 25 °C.

Table 1 .Figure 3 .

 13 Figure 3. Chloride ion conductivity dependence on temperature at full humidification.

Figure 4 .

 4 Figure 4. Chloride ion conductivity (including error bars) as function of the alkaline treatment time.

Figure 6 .

 6 Figure 6. Water uptake as function of the alkaline treatment time. a) PPO-DABCO (OH and Cl form). b) PPO-DABCO + LDH (OH vs Cl form).

Figure 7 .Figure 8 .

 78 Figure 7. Hydration number of composite PPO-DABCO + LDH as function of the alkaline treatment time.

Figure 8a shows a

  Figure8ashows a comparison of typical thermogravimetric curves for pristine PPO-DABCO and PPO-DABCO+LDH after alkaline degradation. For PPO-DABCO, the thermogram shows several, well-separated mass losses: after the initial loss of water (5 %), a broad mass loss above 200 °C may be attributed to the removal of DABCO, which was previously

1 Figure 9 .

 19 Figure 9. FTIR spectra of pristine PPO-DABCO+LDH (black line) and aged PPO-DABCO +LDH in Cl form (red line).

Figure 9

 9 Figure9shows the FTIR analysis of PPO-DABCO+LDH in chloride form before and after the treatment in 2 M KOH at 60°C for 168 h. Absorptions of the PPO backbone are present in both spectra: the stretching vibration of phenyl rings at 1600 and 1470 cm -1 and the

  This large band is due to vibrations of adsorbed water, vibration of HO • • • OH moieties, and vibrations of structural OH groups present in the LDH structure. The absorption maximum is at 3365 cm -1 and 3410 cm -1 for pristine and aged membranes respectively. The OH stretch is generally affected by the strength of the hydrogen bond that depends on the surrounding environment; the strongest hydrogen bonds absorb at the lowest vibrational frequency [51]. Typical C-N vibrations are observed for both samples at 1020 cm -1 [52] that overlap with the deformation mode of Al-OH [53]. Quaternary ammonium moieties absorb at 1060 and 1470 cm -1 [54]. The latter signal overlaps with PPO modes, resulting in a single broad band in the spectrum. The band at 1060 cm -1 shows a lower intensity in the aged membrane indicating some loss of ammonium moieties although the absorption at 1370 cm -1 , due to out-of-plane bending of CH 2 and typical of DABCO derivative, is unchanged [37]. The high intensity of this band in the aged composite is probably due to the overlap of the absorption of CO 3 2-ions, reported around 1362 cm -1 , present in the LDH interlayer after the long treatment in KOH [55]. The main differences between pristine and aged samples are in the fingerprint region especially in the range 500-650 cm -1 where the two typical signals of DABCO, centred at 550 and 610 cm -1 appear, indicating a modification in the aged membrane [56]. The band at 550 cm -1 also overlaps with hydroxyl absorption due to the Mg-Al LDH [53]. Very similar spectra are observed for PPO-DABCO samples before and after the aging.

Acknowledgment

The authors want to thank Prof. F. Costantino (Univ. Perugia, Italy) for recording X-ray diffractograms of the LDH powders.

Conflict of Interest

The authors declare no conflict of interest.


Ping Fan

Z.-K Fan

F Luo

Xianghua Wang

Hongli Zhang

B Ma

Xue

Yang-Hai Xu

Wen-Zhi Li

Bo Fan

Zhong-Kuan Luo

Fang Wang

Xiang-Hua Zhang

Hong-Li Ma

Bai Xue
email: baixue@szu.edu.cn.

Stabilizing electrode/electrolyte interface in Li-S batteries using liquid/solid Li 2 S-P 2 S 5 hybrid electrolyte

Keywords: lithium thiophosphates, 75Li 2 S-25P 2 S 5, hybrid electrolyte, Li-S battery, interfacial stability

Stabilizing electrode/electrolyte interface in Li-S batteries using liquid/solid Li2S-P2S5 hybrid electrolyte

Introduction

Lithium-sulfur batteries, possessing a theoretical specific capacity of 1672 mAh/g and an average voltage of 2.15 V [START_REF] Manthiram | Rechargeable Lithium-Sulfur Batteries[END_REF][START_REF] Ji | A highly ordered nanostructured carbon-sulphur cathode for lithium-sulphur batteries[END_REF], is promising to construct energy storage system for electrical vehicles running more than 400 km [START_REF] Bruce | Li-O 2 and Li-S batteries with high energy storage[END_REF]. However, it encounters serious challenges including poor conductivity of sulfur and lithium sulfides, shuttle problems of lithium polysulfides and uncontrolled growth of lithium dendrites, which result in poor rate capability [START_REF] Cheon | Rechargeable Lithium Sulfur Battery[END_REF], capacity fading [START_REF] Ji | Advances in Li-S batteries[END_REF], coulombic efficiency declining and even cell short-circuit [START_REF] Hayashi | Development of Sulfide Solid electrolytes and interface Formation Processes for Bulk-Type All-Solid-State Li and Na Batteries[END_REF]. To cope with polysulfides shuttling, various strategies are focused on modifying cathode structure [START_REF] Fan | Advanced chemical strategies for lithium-sulfur batteries: A review[END_REF][START_REF] Zheng | Amphiphilic surface modification of hollow carbon nanofibers for improved cycle life of lithium sulfur batteries[END_REF][START_REF] Zhang | Enhancement of long stability of sulfur cathode by encapsulating sulfur into micropores of carbon spheres[END_REF], tuning electrolyte composition [START_REF] Shim | The Lithium/Sulfur Rechargeable Cell: Effects of Electrode Composition and Solvent on Cell Performance[END_REF][START_REF] Sheng | Role of LiNO 3 in rechargeable lithium/sulfur battery[END_REF][START_REF] Adams | Long term stability of Li-S batteries using high concentration lithium nitrate electrolytes[END_REF], optimizing the interface [START_REF] Choi | Effects of carbon coating on the electrochemical properties of sulfur cathode for lithium/sulfur cell[END_REF][START_REF] Fan | Interlayer Material Selection for Lithium-Sulfur Batteries[END_REF] and controlling the electrochemical process [START_REF] Yuan | New insight into the discharge process of sulfur cathode by electrochemical impedance spectroscopy[END_REF][START_REF] Yang | A separator-free lithium/polysulfide semi-liquid battery for large-scale energy storage[END_REF]. However, the shuttle problems cannot be completely eliminated.

Replacing the conventional organic liquid electrolyte with polymer electrolytes [START_REF] Lin | Lithium-sulfur batteries: from liquid to solid cells[END_REF] or dense inorganic solid electrolytes [START_REF] Hayashi | All-solid-state Li/S batteries with highly conductive glass-ceramic electrolytes[END_REF][START_REF] Agostini | A lithium-sulfur battery using a solid, glass-type P 2 S 5 -Li 2 S electrolyte[END_REF], provides an opportunity to eliminate the polysulfides shuttling problem. Simply utilizing polymer electrolytes [START_REF] Huang | Polymers for High Performance Li-S Batteries: Material Selection and Structure Design[END_REF] or solid electrolytes usually leads to unsatisfactory rate performance of the batteries due to the poor ionic conductivity at room temperature and poor contact and instability issues between electrolytes and electrodes [START_REF] Xu | Interfacial challenges and progress for inorganic all-solid-state lithium batteries[END_REF][START_REF] Sun | Recent advances in all-solid-state rechargeable lithium batteries[END_REF]. As a result, it gives birth of the concept of hybrid Li-S batteries, where inorganic solid electrolyte being used as polysulfide barrier and organic electrolyte [START_REF] Wang | A shuttle effect free lithium sulfur battery based on a hybrid electrolyte[END_REF][START_REF] Keller | Hybrid electrolytes for lithium metal batteries[END_REF], ionic liquid electrolyte [START_REF] Lu | Ionic-liquid-nanoparticle hybrid electrolytes: applications in lithium metal batteries[END_REF][START_REF] Sheng | Ionic conductivity promotion of polymer electrolyte with ionic liquid grafted oxides for all-solid-state lithium-sulfur batteries[END_REF] or gel-polymers [START_REF] Wang | A gel-ceramic multi-layer electrolyte for long-life lithium sulfur batteries[END_REF][START_REF] Yu | Hybrid Lithium-Sulfur Batteries with a Solid Electrolyte Separator and Lithium Polysulfide Catholyte[END_REF] used as wetting agent to improve interfacial contact and as high-speed bridge for fast Li + transportation [START_REF] Judez | Polymer-Rich Composite Electrolytes for All-Solid-State Li-S Cells[END_REF][START_REF] Yu | Electrode-Electrolyte Interfaces in Lithium-Sulfur Batteries with Liquid or Inorganic Solid Electrolytes[END_REF]. Oxides, such as LAGP (Li 1.5 Al 0.5 Ge 1.5 (PO 4 ) 3 ) [START_REF] Keller | Hybrid electrolytes for lithium metal batteries[END_REF][START_REF] Gu | A hybrid electrolyte for long-life semi-solid-state lithium sulfur batteries[END_REF] and LATP (Li 1.3 Al 0.3 Ti 1.7 (PO 4 ) 3 ) [START_REF] Wang | Durability of the Li 1+x Ti 2-x Al x (PO 4 ) 3 solid electrolyte in lithium-sulfur batteries[END_REF], etc., are the most used solid electrolytes for inorganic separators. However, their low ionic conductivity and poor contact with cathode cannot satisfy high rate performance. Sulfide solid electrolytes (SSEs), such as, Li 10 GeP 2 S 12 (12 mS cm -1 ) [START_REF] Kamaya | A lithium superionic conductor[END_REF], Li 7 P 3 S 11 (17 mS cm -1 ) [START_REF] Seino | A sulphide lithium super ion conductor is superior to liquid ion conductors for use in rechargeable batteries[END_REF], Li 10 SnP 2 S 12 (4 mS cm -1 ) [START_REF] Bron | Li 10 SnP 2 S 12 : An Affordable Lithium Superionic Conductor[END_REF], Li 9.54 Si 1.74 P 1.44 S 11.7 Cl 0.3 (25 mS cm -1 ) [START_REF] Kato | High-power all-solid-state batteries using sulfide superionic conductors[END_REF] and Li 6 PS 5 Cl (3.15 mS cm -1 ) [START_REF] Wang | High-Conductivity Argyrodite Li 6 PS 5 Cl Solid Electrolytes Prepared via Optimized Sintering Processes for All-Solid-State Lithium-Sulfur Batteries[END_REF], are more attractive than oxides due to their superior ionic conductivity and feasible deformability at ambient temperature.

Despite such advantages, many highly-conductive sulfide electrolytes have an unstable interface against Li anode [START_REF] Wenzel | Interphase formation and degradation of charge transfer kinetics between a lithium metal anode and highly crystalline Li 7 P 3 S 11 solid electrolyte[END_REF][START_REF] Ong | Phase Stability, Electrochemical Stability and Ionic Conductivity of The Li 10±1 MP 2 X 12 (M = Ge, Si, Sn, Al or P, and X = O, S or Se) Family of Superionic Conductors[END_REF]. For example, the reaction between Li 10 GeP 2 S 12 and Li leads to the formation of a continuously growing interphase layer composed of Li 3 P, Li 2 S, and Li-Ge alloy, which increases the interfacial resistance [START_REF] Wenzel | Direct Observation of the Interfacial Instability of the Fast Ionic Conductor Li 10 GeP 2 S 12 at the Lithium Metal Anode[END_REF]. In order to stabilize Li-sulfide electrolytes interface, a solid electrolyte interphase (SEI) is necessarily built up between them [START_REF] Xu | Interface engineering of sulfide electrolytes for all-solid-state lithium batteries[END_REF]. LiF is proposed as one key component in SEI due to its excellent electronic insulating and ionic conducting property, not only allowing sufficiently fast Li-ion kinetics but also more importantly promoting the electrodeposition of Li in a parallel rather than vertical manner [START_REF] Lu | Stable lithium electrodeposition in liquid and nanoporous solid electrolytes[END_REF][START_REF] Yuan | Regulating Li deposition by constructing LiF-rich host for dendrite-free lithium metal anode[END_REF].

In situ modification of the interface between Li metal and Li 10 SnP 2 S 12 with specific ionic liquid and salts leads to the enrichment of LiF in the SEI layer [START_REF] Zheng | Stabilizing Li 10 SnP 2 S 12 /Li Interface via an in Situ Formed Solid Electrolyte Interphase Layer[END_REF]. An in situ and self-limiting reaction between P 4 S 16 and Li in N-methyl-2-pyrrolidone enables formation of Li 3 PS 4 solid-state electrolyte protection layer on the surface of Li, which can also address the dynamic plating/stripping process of Li metal [START_REF] Liang | In Situ Li 3 PS 4 Solid-State Electrolyte Protection Layers for Superior Long-Life and High-Rate Lithium-Metal Anodes[END_REF].

The stability between solid electrolytes and cathodes is also a key factor that determines the performance of hybrid batteries. For example, in hybrid Li-S batteries the NASICON-type solid electrolyte Li 1+x Ti 2-x Al x (PO 4 ) 3 gets lithiated in contact with lithium polysulfide solution and Li 2 CO 3 is formed on the particle surface, blocking the interfacial lithium-ion transport between the liquid and solid-state electrolytes [START_REF] Wang | Durability of the Li 1+x Ti 2-x Al x (PO 4 ) 3 solid electrolyte in lithium-sulfur batteries[END_REF]. In our recent work, the hybrid Li-S batteries, based on Li 10 SnP 2 S 12 sulfide solid electrolyte, successfully run more than 100 cycles under an extremely high rate of 5C (7.10 mA/cm 2 ) and maintain good capacity and coulombic efficiency [START_REF] Xu | An extremely high rate Li-S battery with hybrid electrolyte[END_REF]. However, Li 10 SnP 2 S 12 , which contains tetravalent Sn 4+ , is not so stable with both the cathode and the anode, thus initiating continuously unwanted parasitic reactions.

Here, lithium thiophosphates xLi 2 S•(100-x)P 2 S 5 (70≤x≤75, SSEs), which exclude tetravalent metal elements, are used as the solid electrolyte separator for hybrid Li-S batteries. The influence of Li 2 S content of lithium thiophosphates on their interfacial stability with the Li anode and the cathode is detailly investigated. After achieving both stable interfaces with the Li anode and the cathode by adjusting the Li 2 S content, the lithium thiophosphates 75Li 2 S•25P 2 S 5 is used to fabricate a high-performance Li-S battery which successfully runs more than 400 cycles at 0.2C.

Experimental

Materials preparation

Lithium thiophosphates xLi 2 S•(100-x)P 2 S 5 (70≤x≤75) were prepared by conventional ball-milling method plus low temperature annealing. 4 g mixture of Li 2 S (99.99%, Alfa Aesar) and P 2 S 5 (Sigma-Aldrich, ≥99%) was put into a 45-ml tungsten carbide jar with three tungsten carbide grinding balls (φ=10mm). The content of Li 2 S (mole ratio) was precisely controlled to 70%, 71.25%, 72.5%, 73.75% and 75%, respectively. Then ball milling was conducted at 450 rpm in intermittent manner leads to sulfur evaporation. The effective milling time was 70 h. To gain better homogeneity, the agglomerated powder was intermittently ground every 8-h milling.

All the operations were carried out in an argon-filled glove-box (O 2 <0.5 ppm, H 2 O< 0.5 ppm). The resultant Li 2 S-P 2 S 5 glass powder was then cold-pressed into pellets (φ15 mm × 1.1 mm) under a pressure of 330 MPa. The pellets were sealed in silica glass tubes under vacuum (10 -3 Pa) and sintered at 260 °C for 1 h. The highly ionic conductive lithium thiophosphate separators (SSEs) were then obtained.

Materials characterization method

X-ray diffraction (D8 Advance, Bruker AXS GmbH, Germany) and Raman spectra (inVia, Renishaw Inc., UK; loaded with a 532 nm diode-pumped solid-state laser as the excitation source) were used to monitor the phase evolution during the sample preparation. To avoid moisture, they were mounted in the sample holder and sealed by polyimide films in the glove-box prior to XRD and Raman characterization.

The chemical stability of as-obtained sulfide electrolytes with liquid electrolyte was checked by soaking experiments. 300 mg powder of lithium thiophosphates xLi 2 S•(100-x)P 2 S 5 (70≤x≤75) was firstly immersed in 5 ml DME-DOL solvent (50-50,v%) for 48 h. The mixture was vigorously shaken every 8 h and finally vacuum dried at 80 o C for composition analysis by XRD and Raman spectra. The morphology of the samples after battery test was recorded by a field-emission scanning electron microscope (SU8010, Hitachi Inc., Japan), which was equipped with an energy dispersive X-ray micro-analyzer for analyzing the composition of the samples. X-ray photoelectron spectroscopy (XPS) of the samples after battery test was measured using Thermo Scientific ESCALAB 250Xi with Al Kα-radiation. Vacuum transfer box was used to transfer the samples from glove-box to the characterization equipment. Li foils for SEM and XPS analysis were rinsed by dimethoxyethane (DME) three times in glove-box, and dried under vacuum for one night before test.

Thermal analysis of the lithium thiophosphates was carried out by using a differential scanning calorimeter (DSC 3, Mettler Toledo Inc., Switzerland) with a scanning rate of 10 °C/min under N 2 flow.

The ionic conductivity was measured using an AC impedance method in an argon filled glove-box with a frequency response analyzer (Solartron 1260A, Solartron Analytical Inc., UK). The as-obtained lithium thiophosphate separators were deposited with gold on both two sides before test. The conductivity of the samples was then determined from the impedance spectra data. The electronic conductivities of lithium thiophosphates were measured by using a DC polarization method with a constant DC voltage of 0.1 V working on the Au/SSEs/Au cells [START_REF] Han | High electronic conductivity as the origin of lithium dendrite formation within solid electrolytes[END_REF] on the CHI660E electrochemical workstation (CHI660E, CH Instruments Ins., Shanghai, China).

Electrochemical measurement

Button type hybrid Li-S batteries (CR2032) were assembled in a argon-filled glove-box (O 2 < 0.5 ppm, H 2 O < 0.5 ppm) with a structure of -)SUS/Li/ liquid electrolyte (LE)/solid electrolyte (SSEs)/liquid electrolyte/S-C/SUS(+. The cathode was obtained by ultrasonic spraying the ethanol suspension of S-C composite, LA133 binder and conductive acetylene black (8:1:1, wt%) on carbon paper, where the S-C composite was prepared in advance by heating the ground mixture of ketjen black and S at 155 °C for 12 h under vacuum with 60 wt% sulfur loading in the composite. The sulfur load in the cathode was 0.85 mg cm -2 . The cathode and both sides of the SE separator were wetted by organic liquid electrolyte (10mL (Anode)/ 30mL (Cathode) LE, 1 M LiTFSI DME/DOL (dimethoxyethane/1,3-dioxolane, 50-50 vol%)-1% LiNO 3 ). The batteries were stored in the glove-box for one night before galvanostatic charge-discharge test on a battery tester (CT2001, Land Ins., Wuhan, China). The cutoff voltage was 1.5 and 2.8 V. The impedance spectra of the batteries were recorded by an electrochemical workstation.

Results and discussion

Materials characterization

Lithium thiophosphates were synthesized using high energy ball milling plus low temperature annealing. thiophosphates at 406 cm -1 , 422 cm -1 and 388 cm -1 can be attributed to P 2 S 7 4-, PS 4 3- and P 2 S 6 4-, respectively [START_REF] Wei | Influence of annealing on ionic transfer and storage stability of Li 2 S-P 2 S 5 solid electrolyte[END_REF][START_REF] Calpa | Instantaneous preparation of high lithium-ion conducting sulfide solid electrolyte Li 7 P 3 S 11 by a liquid phase process[END_REF]. There's a tendency that the signal of P 2 S 7 4-, which is a principal structural unit of Li 7 P 3 S 11 , weakens as Li 2 S content increase from x=70 to x=75, while the signal of PS 4 3-strengthens synchronously. This is in good agreement with the previous XRD results, which show that the major phase of the glass-ceramic gradually transfers from Li 7 P 3 S 11 to Li 3 PS 4 when the Li 2 S content increases from 70% to 75%. The signal of P 2 S 6 4-is always observed and maintains the same level in the Raman spectroscopy. The signal of P 2 S 6 4-unit is less intense in the corresponding as-milled samples than the annealed samples (Figure S2), reflecting the byproducts are mainly generated during the annealing [START_REF] Kudu | A review of structural properties and synthesis methods of solid electrolyte materials in the Li 2 S -P 2 S 5 binary system[END_REF].

Figure 1c

shows the DSC curves of the as-milled amorphous lithium thiophosphates. It's observed that there are two crystallization peaks for all the samples, which is in accordance with that reported by Zhang [START_REF] Zhang | Synergistic effect of processing and composition x on conductivity of xLi 2 S-(100-x)P 2 S 5 electrolytes[END_REF]. For the sample x=70, the two crystallization peaks are very close and positioned at 273.5 o C (T1, low crystallization temperature) and 276.9 o C (T2, high crystallization temperature). With the increase of Li 2 S content, both of the two crystallization peaks shift to lower temperature (Table 1) and the two crystallization peaks are separated. Meanwhile, the intensities of the T1 peak get strengthened while the intensities of the T2 peak get weakened. The two crystallization peaks at T1 and T2 are probably assigned to the formation of Li 3 PS 4 and Li 7 P 3 S 11 , respectively. The DSC measurement is a dynamic process during which the temperature rises fast (10 o C /min), and usually shifts the curves to higher temperature than isothermal processes. As shown in Figure 1c, there's obvious onset signal for exothermic process around 260 o C, that is, the crystallization starts. It's feasible to complete the crystallization by extending the annealing time at the onset crystallization temperature. Therefore, the annealing temperature of lithium thiophosphates is chosen to be 260 o C to get sufficient crystallization [START_REF] Mizuno | High lithium ion conducting glass-ceramics in the system Li 2 S-P 2 S 5[END_REF] and to avoid the undesired decomposition of Li 7 P 3 S 11 into Li 4 P 2 S 6 at higher temperature which is detrimental to the conductivity [START_REF] Wei | Influence of annealing on ionic transfer and storage stability of Li 2 S-P 2 S 5 solid electrolyte[END_REF]. The annealed sample of x=75 exhibits pure phase of Li 3 PS 4 similar to x=72.5 and x=73.75 but lower crystallization temperature, which is probably associated with its small grain size and fast nucleation rate mentioned above. This work will not further discuss the influence of the temperature change on the crystallinity, conductivity and structure of SSEs, but just focuses on one temperature of 260 o C. R is resistance determinate by the arc intercept at low frequency. The DC polarization method as described in Ref [START_REF] Han | High electronic conductivity as the origin of lithium dendrite formation within solid electrolytes[END_REF] is used to estimate the electronic conductivity. The information about the ionic conductivity and electronic conductivity of SSEs are summarized in Table S1.

The ionic conductivity of the as-milled lithium thiophosphates firstly increases from 0.16 mS/cm at x=70 to 0.40 mS/cm at x=73.75 and then decreases a little. The corresponding annealed samples show a contrary trend, where the ionic conductivity exhibits the maximum of 3.3 mS/cm at x=70 and then gradually decreases to about 0.40 mS/cm. The superior ionic conductivity of the annealed sample at x=70 is related with the formation of highly ionic conductive phase Li 7 P 3 S 11 , and its decrease with the increase of Li 2 S content is due to the transition of the crystalline phase from especially when the solid electrolyte directly contacts with lithium anode. This drawback is expected to be compensated by introducing a small amount of organic liquid electrolyte between the solid electrolyte and the lithium anode, which will be demonstrated as below in our hybrid Li-S batteries with the lithium thiophosphate separators . Coulombic efficiency of all the hybrid Li-S batteries keeps at about 100% over the whole cycles (Figure 2f). It is therefore confirmed that regardless of their stability issues, all the lithium thiophosphate separators with the studied compositions can effectively block the diffusion of polysulfide ions from cathode to anode. The polysulfide shuttling is no longer a main cause of performance degradation in these hybrid batteries. The distinct difference of the battery performance is principally originated from the different chemical/electrochemical stability of the lithium thiophosphate separators with different Li 2 S content, which will be discussed in the following three sections.

Battery performance

Chemical stability with Ether-Based Solvent

To evaluate the influence of the composition of the lithium thiophosphates on the

Electrochemical stability with Li anode

A stable interface between electrolyte and lithium anode is essential for a lithium The P 2p spectrum exhibits several peaks, which can be fitted with three distinct doublets (2p1/2 and 2p3/2) (Figure 5g,5k). Among them, the P 2p3/2 peaks appear at 131.7 eV and 132.3 eV can be assigned to PS 4 3-[45,61]. This assignment is further confirmed by the S 2p spectrum with one intensive responsive peak of S 2p3/2 at 161.6 eV attributed to PS 4 3-(Figure 5h, 5l) [START_REF] Liang | In Situ Li 3 PS 4 Solid-State Electrolyte Protection Layers for Superior Long-Life and High-Rate Lithium-Metal Anodes[END_REF][START_REF] Dietrich | Spectroscopic characterization of lithium thiophosphates by XPS and XAS -a model to help monitor interfacial reactions in all-solid-state batteries[END_REF][START_REF] Zu | Activated Li 2 S as a High-Performance Cathode for Rechargeable Lithium-Sulfur Batteries[END_REF]. Assigning two different binding energies to PS 4 3-reflects that this anion locates in environment with different crystallinity or structural disorder, since amorphization often shifts binding energies in XPS spectra [START_REF] Tanibata | A novel discharge-charge mechanism of a S-P 2 S 5 composite electrode without electrolytes in all solid-state Li/S batteries[END_REF]. The strong signal of PS 4 3-reveals the existence of Li 3 PS 4 in the SEI layer on Li anode. The P 2p3/2 peak positioned at 133.2 eV is very weak which is probably related with P-S-P bridging sulfur structure and reflects its low content on Li anode [START_REF] Dietrich | Spectroscopic characterization of lithium thiophosphates by XPS and XAS -a model to help monitor interfacial reactions in all-solid-state batteries[END_REF]. The S 2p3/2 signal also denotes the existence of -S * O 2 CF 3 at 169.3 eV and -S * O 2 CF 2 + at 167.2 eV from the decomposition products of LiTFSI [START_REF] Busche | Dynamic formation of a solid-liquid electrolyte interphase and its consequences for hybrid-battery concepts[END_REF], and a trace of Li 2 S at 160.0 eV [START_REF] Aurbach | On the Surface Chemical Aspects of Very High Energy Density Rechargeable Li-Sulfur Batteries[END_REF]. F 1s Signals located at 684.9 eV and 685.2 eV belonging to LiF are also detected (Figure 5i and Figure 5m), which are probably related to the decomposition product of LiTFSI salt with Li metal [START_REF] Busche | Dynamic formation of a solid-liquid electrolyte interphase and its consequences for hybrid-battery concepts[END_REF][START_REF] Fan | Fluorinated solid electrolyte interphase enables highly reversible solid-state Li metal battery[END_REF]. A LiF-rich SEI (above 9%, Figure S7, Table 2) is helpful to suppress the penetration of Li dendrites into SSEs, while the low electronic conductivity and the intrinsic electrochemical stability of LiF block side reactions between the SSEs and Li [START_REF] Lu | Stable lithium electrodeposition in liquid and nanoporous solid electrolytes[END_REF][START_REF] Fan | Fluorinated solid electrolyte interphase enables highly reversible solid-state Li metal battery[END_REF].Two Li 1s peaks appeared at 54.8 eV and 55.7 eV (Figure 5j, 5n) can be assigned to Li-S bonding [START_REF] Liang | In Situ Li 3 PS 4 Solid-State Electrolyte Protection Layers for Superior Long-Life and High-Rate Lithium-Metal Anodes[END_REF][START_REF] Kim | Lithium-Sulfur Batteries: Stabilization of Insoluble Discharge Products by Facile Aniline Modification for High Performance Li-S Batteries[END_REF] and Li-O bonding, respectively [START_REF] Chu | Cobalt ferrite thin films as anode material for lithium ion batteries[END_REF].

As shown in

Electrochemical stability with cathode

The interfacial stability between SSEs and cathode attracts less attention of the electrochemical community than that with lithium anode, but has been unambiguously clarified to be essential for the cycling performance of sulfide-liquid electrolyte-based hybrid Li-S batteries by our previous study [START_REF] Xu | An extremely high rate Li-S battery with hybrid electrolyte[END_REF]. This section will focus on the electrochemistry stability of the lithium thiophosphates with the cathode wetted by ether-based liquid electrolyte. This is firstly characterized by utilizing cyclic voltammetry (CV) method in an asymmetric cell with structure of -)

Li/LE-SSEs/LE/SS (stainless-steel)(+ within a potential range of 0.

  To study the electrochemical stability of different solid electrolyte separators in solid-liquid-hybrid environment, Cyclic voltammetry of -)Li/liquid electrolyte (LE)/solid electrolyte (SSEs)/liquid electrolyte (LE)/SUS(+ asymmetric cells was measured on the CHI660E electrochemical workstation. 10 μl liquid electrolyte (1 M LiTFSI DME/DOL (50-50 vol%)-1% LiNO 3 ) was dropwise added between the SE separator and the stainless-steel electrode to create the solid-liquid electrolyte interface. Symmetric cells with a structure of Li/liquid electrolyte (LE) /solid electrolyte (SSEs) /liquid electrolyte (LE) /Li were assembled in Swagelok battery moulds. 10 μl liquid electrolyte was dropped between the SSEs separator and the Li electrode. Li plating/stripping behaviours of the symmetric cells were measured by the CT2001 battery tester.

  Figure 1a and Figure S1 shows the X-ray diffraction (XRD) spectra of the glass-ceramics and glass samples in the composition range 70≤x≤75 for xLi 2 S•(100-x)P 2 S 5 . No diffraction peaks are observed for the glass samples after 70 h of ball milling, indicating their amorphous state. After annealing at a low temperature of 260 o C for 1 h, all the glass samples transform into glass-ceramics. The diffraction peaks match well with the Li 7 P 3 S 11 crystalline at x=70 [48]. As the increase of Li 2 S content, the major crystalline phase transforms into Li 3 PS 4 gradually. There exist two phases of Li 7 P 3 S 11 and Li 3 PS 4 at x=71.25, which transforms into Li 3 PS 4 after x≥72.5. The glass-ceramic shows integrated diffraction peaks attributed to Li 3 PS 4 at x=75. It has broad diffraction peaks compared with that of x=72.5 and x=73.75, which is possibly related with its smaller grain size originating from higher driving force of nucleation during crystallization. As shown in Figure 1b and Figure S2, the local structure of lithium thiophosphates xLi 2 S•(100-x)P 2 S 5 (70≤x≤75), in both glass-ceramic and glass forms was examined by Raman spectroscopy. The three bands in annealed lithium

Figure 1 .

 1 Figure 1. The XRD patterns, Raman spectra, DSC curves and Ionic conductivities of the lithium

Figure

  Figure 1dexhibits the room temperature ionic conductivity of the lithium

4 .

 4 Figure S5. The capacity of the batteries decreases during cycling for all the cases

Figure 2 .

 2 Figure 2. Galvanostatic charge-discharge profiles of the hybrid Li-S batteries at 0.2 C. (a) The

  electrochemical performance of the hybrid Li-S batteries, the chemical stability of the as-obtained sulfide electrolytes with liquid electrolyte is nonnegligible and should be checked first. The soaking experiment was carried out by immersing 300 mg of lithium thiophosphates xLi 2 S•(100-x)P 2 S 5 (70≤x≤75) into 5ml DME-DOL solvent (50-50,v%) for 48 h. The mixture was vigorously shaken every 8 h and finally vacuum dried at 80 o C for composition analysis by XRD and Raman spectra. As shown by the XRD patterns in the Raman spectra (Figure3b), all the five soaked SSEs display almost identical diffraction peaks which can be attributed to a crystalline phase DME-solvated Li 3 PS 4[START_REF] Ito | Synthesis of Crystalline Li 7 P 3 S 11 Solid Electrolyte from 1,2-Dimethoxyethane Solvent[END_REF][START_REF] Xue | Solvent-Assisted Ball Milling for Synthesizing Solid Electrolyte Li 7 P 3 S 11[END_REF]. There appear two distinctive peaks positioned at 421cm -1 and 388 cm -1 in Figure3b, corresponding to ether-solvated PS 4 3-anion and ether-solvated vertex-shared PS 4 tetrahedral anions[START_REF] Ito | Synthesis of Crystalline Li 7 P 3 S 11 Solid Electrolyte from 1,2-Dimethoxyethane Solvent[END_REF], respectively. Similar Raman signals have been identified in recent studies on liquid-phase synthesis of Li 7 P 3 S 11 in acetonitrile (ACN), where the high wavenumber signal is attributed to a crystalline phase Li 3 PS 4 •ACN[START_REF] Wang | Mechanism of Formation of Li 7 P 3 S 11 Solid Electrolytes through Liquid Phase Synthesis[END_REF] and the low wavenumber signal is attributed to an amorphous phase "Li 2 S•P 2 S 5 "[START_REF] Fan | Influence of precipitate/supernatant ratio during liquid-phase synthesis of solid electrolyte Li 7 P 3 S 11[END_REF][START_REF] Fan | Will Sulfide Electrolytes be Suitable Candidates for Constructing a Stable Solid/Liquid Electrolyte Interface?[END_REF]. Therefore, it is supposed that the lithium thiophosphates are decomposed into these two phases by the solvent in our case. It is worth noting that the intensity of the signal at 388 cm -1 decreases with the increasing Li 2 S content, indicating that less "Li 2 S•P 2 S 5 " releases, especially for the SSEs x=73.75 and x=75.

Figure 3 .

 3 Figure 3. Chemical stability of Li 2 S-P 2 S 5 -based glass-ceramic SEs with ether-based solvent. (a)

  metal battery. Li plating/stripping behaviors of Li-LE-SSEs-LE-Li symmetric cells, as shown in Figure4a, are characterized to explore the interfacial stability of the lithium thiophosphate separators under reduction condition. All the Li symmetric cells, excluding the case of x=71.25 and x=72.5, display a steady voltage below 0.08 V and capacity of 0.99 mAh under 0.23 mA cm -2 for a long term of more than 200 h, revealing an excellently stable interface between lithium thiophosphate separators and lithium anode with the existence of liquid electrolyte. The Li electrodeposition of 0.99 mAh at a current density of 0.23 mA cm -2 is competitive in solid-state Li symmetric cells[START_REF] Shao | Drawing a Soft Interface: An Effective Interfacial Modification Strategy for Garnet-Type Solid-State Li Batteries[END_REF]. It's puzzled for the short life span of x=71.25 and x=72.5. Repeated tests of the Li/SSEs/Li symmetric cells were conducted, however short circuit still occurred unavoidably and randomly even for the SSEs x=73.75 with the best cycling performance (840 h vs 315 h, see FigureS6). Cracks accidentally formed during cold pressing and sintering of the separators can cause lithium penetration and short circuit of the batteries[START_REF] Nagao | In situ SEM study of a lithium deposition and dissolution mechanism in a bulk-type solid-state cell with a Li 2 S-P 2 S 5 solid electrolyte[END_REF]. The pores and the flaws cannot be avoided in the SSEs pellets which are simply prepared by cold pressing. Furthermore, dendrites can also form at the grain boundaries or even in amorphous sulfide electrolytes free of grain boundaries[START_REF] Porz | Mechanism of Lithium Metal Penetration through Inorganic Solid Electrolytes[END_REF]. The discussion of lithium penetration is beyond the scope of this article. However, from the perspective of stability, these two samples both exhibit stable voltages before short-circuited, indicating their good interfacial stability with lithium anode like other compositions.The electrochemical impedance spectroscopy (EIS) of the Li symmetric cells is tested and recorded in Figure4b-c. Two arcs are observed in the EIS for all the Li symmetric cells before and after 30 h Li plating/stripping test. The first arc at high frequency could be attributed to the interfacial resistance, which is impressively decreased after 30 h continuously plating/striping. The Li symmetric cells for lower Li 2 S content including x=70, x=71.25 and x=72.5 exhibit small bulk resistance (matches well with Figure1d) but large interfacial resistance and charge transfer resistance. Generally, the total resistance of the three samples is small, which corresponds to the low overpotential of the Li symmetric cells that is below 0.05 V.On the contrary, Li symmetric cells of x=73.75 and x=75 show large bulk resistance but small interfacial resistance and charge transfer resistance, therefore leading to relatively high overpotential of the symmetric cells. The larger interfacial resistance for cells containing lower Li 2 S content is believed to be related with the formation of thick SEI layer, which will be further demonstrated by SEM and XPS later.

Figure 4 .

 4 Figure 4. Stability of sulfide solid electrolyte (SSEs)/liquid electrolyte (LE) interface with the

Figure 5

 5 Figure 5 displays the morphology (Figure 5a-f) and composition changes

Figure 5 .

 5 Figure 5. SEM images and XPS spectra of the Li metal surface after battery test. (a-f)

7 V

 7 3 -3.5 V, which covers the operating voltage of the Li-S battery cathode. As shown in Figure 6a, the CV curves show an onset oxidation current at about 2.5 V which dramatically increases to 67 μA at 3.5 V for x=70. During the subsequent cathodic sweep, an onset reduction current occurs at about 2.3 V and transforms into sharp reduction current peak at 1.5 V. There's similar redox potential for the SSEs of x=71.25 and x=72.5however with relatively lower anodic and cathodic current. The anodic/ cathodic current is normally accountable for delithiathon/ lithiation of lithium thiophosphate[START_REF] Zhu | Origin of Outstanding Stability in the Lithium Solid Electrolyte Materials: Insights from Thermodynamic Analyses Based on First-Principles Calculations[END_REF][START_REF] Li | Electrochemically primed functional redox mediator generator from the decomposition of solid state electrolyte[END_REF]. For the cases of x=73.75 and x=75 (corresponding to Li 3 PS 4 ), two pairs of anodic/cathodic current peaks are observed and located at 2.7 V/2.1 V and 1.0 V/0., respectively, all with very low intensity. CV curves for the following 4 cycles exhibit repeatable redox characteristics (see FigureS9-FigureS13). The redox pair that centered at 2.5 V is in good agreement with previous work on Li[START_REF] Bruce | Li-O 2 and Li-S batteries with high energy storage[END_REF] PS 4 reported by M. Li [72]. The extremely low redox current under CV test indicates that samples for x=73.75 and x=75 (mainly Li 3 PS 4 ) are electrochemical stable under oxidation condition. To further investigate the influence of electrochemical reaction on the interface between cathode and hybrid electrolyte, the cathodes after 40 th charge are analyzed by SEM and XPS. For the batteries with x=70, 71.25 and x=72.5, a gel-like thick layer is observed on the cathode that covers the porous structure of S-C composite (Figure 6b-d), while this layer is absent for the batteries with x=73.75 and x=75 (Figure 6e-f) and the porous morphology of the pristine S-C cathode (Figure S14) is almost unchanged. This is in good accordance with the CV results of the asymmetric half cells, where the redox current of the lithium phosphate samples is much higher at x=70, 71.25 and x=72.5 than that at x=73.75 and x=75 under the operating voltage of the Li-S battery cathode. The XPS spectra collected from the cathode surface after 40 cycles are recorded in Figure 6 (g-h) and Figure S15. In addition to the peaks attributed to Li salts and their decomposition products (SO 2 CF 3 , S 2p3/2, 169.

Li 2 S

 2 content tends to be decomposed by ether-based solvent. The partially soluble decomposition product Li 2 S•P 2 S 5 can diffuse to the active sites on the cathode, and therefore promotes the electrochemical decomposition of the lithium thiophosphates. This can be further confirmed by carefully reexamining the galvanostatic charge-discharge profiles of the batteries with x=70, 71.25 (Figure 2b), where an additional discharge plateau is observed at the beginning of discharge stage and an additional charge plateau at the end of charge stage and consequently exceeding theoretical specific capacity. The gel-like layer can impede Li + diffuse to the active site of S therefore deteriorates electrochemical reaction, resulting in the poor cycling stability at x=70, 71.25 and x=72.5. This gel-like layer is absent at x=73.75 and x=75 due to the high stability of these two solid-state separators with the cathode, and consequently good battery cycling stability is achieved.

Figure 6 . 7 .

 67 Figure 6. Stability of the sulfide solid electrolyte (SSEs)/liquid electrolyte (LE) interface with the

Table 1 .

 1 Crystallization temperature (T1, low crystallization temperature; T2, high crystallization temperature) for as-milled amorphous lithium thiophosphates xLi 2 S•(100-x)P 2 S 5 (70≤x≤75) from

	DSC.
	Temperatur			SSEs (Li 2 S content, %)
	e ( o C)	x=70	X=71.25	X=72.5	X=73.75	X=75
	T1	273.5	261.7	258.5	257.3	243.0
	T2	276.9	275.6	273.8	272.9	258.9

  Li 7 P 3 S 11 to Li 3 PS 4 phase as previously demonstrated by both XRD and Raman spectra. The electronic conductivities of all the SSEs are in the level of 10 -8 S/cm, which is about one order of magnitude higher than that of Li 3 PS 4 (2.2 × 10 -9 S/cm)

reported by C. Wang.

[START_REF] Han | High electronic conductivity as the origin of lithium dendrite formation within solid electrolytes[END_REF]

Such a high electronic conductivity enabling the formation of lithium dendrites is normally detrimental for solid-state Li-metal batteries,

Table 2 ,

 2 the component of S 2s peak attributed to PS 4 3-decreases from 68.84 % (x=70) to 15% (x=75), indicating that the lithium thiophosphate with low Li 2 S content favors the formation of Li 3 PS 4 in the SEI layer. Our recent study and above chemical stability analysis together reveal that Li 7 P 3 S 11 tends to be decomposed by ether solvent into solvated Li 3 PS 4 and "Li 2 S•P 2 S 5 "[START_REF] Fan | Will Sulfide Electrolytes be Suitable Candidates for Constructing a Stable Solid/Liquid Electrolyte Interface?[END_REF]. It is noteworthy that solvated Li 3 PS 4 is just solvated rather than dissolved. "Li 2 S•P 2 S 5 " is partially dissolved in the liquid electrolyte, which can diffuse to the Li surface and react with it to form Li 3 PS 4 . This process is similar to the reported studies on in-situ formed Li 3 PS 4 SEI layers, where phosphor-containing additive is deliberately introduced into the liquid electrolyte to react with Li to form the Li 3 PS 4 SEI layers[START_REF] Lin | Phosphorous Pentasulfide as a Novel Additive for High-Performance Lithium-Sulfur Batteries[END_REF][START_REF] Lu | Pre-modified Li 3 PS 4 based interphase for lithium anode towards high performance Li-S battery[END_REF]. Since the lithium thiophosphate with low Li 2 S content contain more Li 7 P 3 S 11 or its characteristic structural unit P 2 S 7 4-, a thick SEI layer rich of Li 3 PS 4 tends to be formed on the surface of Li anode, as shown by the SEM images in Figure5b-d. This SEI layer composed of Li 3 PS 4 and LiF, as demonstrated in the literatures[START_REF] Liang | In Situ Li 3 PS 4 Solid-State Electrolyte Protection Layers for Superior Long-Life and High-Rate Lithium-Metal Anodes[END_REF][START_REF] Fan | Fluorinated solid electrolyte interphase enables highly reversible solid-state Li metal battery[END_REF], can synergistically stabilize the Li metal so that long-term Li plating/stripping cycles is achieved in the Li-LE-SSEs-LE-Li symmetric cells. In summary, lithium thiophosphates with low Li 2 S content favor the formation of an SEI layer rich of Li 3 PS 4 , so exhibit better electrochemical stability with lithium anode.

Table 2 .

 2 A detailed information of PS 4 3-and LiF content based on fitting peak area from XPS of the Li metal surface after battery test.

	SEI layer		Component content (Area %)
	Component	x=70	x=71.25	x=72.5	x=73.75	x=75

  ether-based solvent[START_REF] Liang | In Situ Li 3 PS 4 Solid-State Electrolyte Protection Layers for Superior Long-Life and High-Rate Lithium-Metal Anodes[END_REF][START_REF] Dietrich | Spectroscopic characterization of lithium thiophosphates by XPS and XAS -a model to help monitor interfacial reactions in all-solid-state batteries[END_REF][START_REF] Tanibata | A novel discharge-charge mechanism of a S-P 2 S 5 composite electrode without electrolytes in all solid-state Li/S batteries[END_REF]. This peak is absent for SSEs with high Li 2 S content (x=73.75 and x=75), which corresponds to the absence of gel-like Therefore, it is proposed that the gel-like layer is formed by the redox product of the lithium phosphate separator. As mentioned above, the lithium phosphate with low

	layer in these samples. Feeble S 2p3/2 peak at 161.6 cm -1 and P 2p3/2 peak at 132.3
	cm -1 , which jointly belong to PS 4 3-[45,61,62], can also be observed. This could be
	originated from a trace of residual of lithium thiophosphates during the cell
	disassembling process (x=73.75 and x=75).
	6 cm -1 ;
	SO 2 CF 2 /LiS x O y , S 2p3/2, 167.4 cm -1 ) [64], S 2p3/2 peak positioned at 164.1 cm -1 and

P 2p3/2 peak centered at 134.3 cm -1 can be detected for SSEs with low Li 2 S content, which can be assigned to P-S/S-S due to parasitic reaction between lithium thiophosphates and

Acknowledments

The work is financially supported by the National Science Foundation of China (NFSC) project (No. 51702216), the Natural Science Foundation of Guangdong Province (No. 2020A1515011430), Shenzhen Science and Technology Foundation (JCYJ20180507182106754), State Key Laboratory of Silicon Materials Visiting Scholar Fund (SKL2020-10) and support from Postdoctoral Science Foundation of China.

Conclusions

Lithium thiophosphates xLi 2 S•(100-x)P 2 S 5 (70≤x≤75, SSEs) were synthesized using high energy ball milling plus low temperature annealing, and were used as solid electrolyte separator to fabricate hybrid Li-S batteries. As Li 2 S content increases from x=70 to x=75, the major phase gradually transforms from Li

Declaration of competing interest

The authors declare no competing financial interests.

Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.apsusc. x0xx00000x.

Hybrid Li-S battery (x=75) yields 400 cycles due to its excellent compatibility with the electrodes.

Declaration of interests

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.