Review—Inorganic Solid State Electrolytes: Insights on Current and Future Scope

The three-dimensional framework structure of NASICON [NAtrium Super Ionic CONductor] was proposed by Hong and Goodenough et al. ^10,11 High ionic conductivity and excellent structural stability are the key fascinating properties of NASICONs. Na_1+x Zr₂P_3−x Si_x O₁₂ [0 ≤ x ≤ 3] was the first NASICON type solid-state electrolyte. The 3D framework of AMM'P₃O₁₂ or AMM'[PO₄]₃ NASICON structure is shown in Fig. 3, where "A" is Li and Na, M is Ti, Ge, Sn, Hf, or Zr & M' is Cr, Al, Ga, Sc, Y, In, or La. In the framework of NASICON, the phosphorus can be partially replaced from S, Si, or As. ^12–14 The 3D NASICON structure is made by the robust connection of MO₆ octahedra and [S, Si, As, P]O₄ tetrahedra by sharing the corner oxygen. Over every two octahedral, three tetrahedral are present in this 3D framework structure ¹⁵ Fig. 3 shows the Li-ion occupies two different sites: in 6-fold coordination octahedrally symmetric site M, and/or 8-fold coordinated site M', located between two columns of M'O6 octahedra. The lithium-ion migrates via hopping between these two M and M' interstitial positions through the bottlenecks. The size of the bottleneck depends on the type of skeleton ions and carrier density at both interstitial positions. Consequently, the structure and ionic conduction of NASICON compounds depend on the lattice parameters of the cell, whose values can be varied by selecting chemicals of different compositions. It has been seen that the materials with long lattice parameters show higher ionic conductivity. Also, the cell of the long lattice parameter increases the size of the bottleneck thereby increasing the lithium-ion conduction. ^14,16–18 Choosing M' as Al³⁺ and Cr³⁺ in AMM'[PO₄]₃ induces more lithium ions at vacant M' sites. At the same time, it also increases the unit cell parameter and therefore bottleneck size, which further increases the ionic conductivity in NASICON type solid-state electrolytes. ¹⁹ Though, due to the formation of the secondary phase of Al³⁺ or Sc³⁺ and large ionic radius mismatch, the substitution level can't be taken beyond 15% (i.e. x = 0.3). To date, Li_1.3Al_0.3Ti_1.7 [PO₄]₃ is the best-reported NASICON lithium-ion conductor with bulk conductivity [σ ≈ 3 × 10⁻³ S cm⁻¹] at room temperature. ²⁰

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The term LISICON stands for Lithium-Super-Ionic-Conductor. Hong and his co-workers started the study on the Li_2+2xZn_1−xGeO₄ family and at 300 °C they reported an ionic conductivity of 1.25 × 10⁻¹ S for Li₁₄Zn [GeO₄]₄. ^2,21,22 Similar to γ-Li₃PO₄, the crystal structure of LISICON is made of a rigid three dimensional [Li₁₁Zn [GeO₄]₄]³⁻ network and three loosely bonded Lithium ions having an orthorhombic unit cell of lattice parameters a = 10.828Å, b = 6.251 Å, c = 5.140 Å, and z = l, and Pnma space group with tetrahedrally co-ordinated cations. ^14,22 Similar to other NASICON like super-ionic conductors, LISICON also contains a high density of mobile lithium ions occupying the sites in the interstitial space of the 3D network. In this network, two factors govern the conduction of Li-ion − 1] the bottleneck size between adjacent sites i.e. 4.38 Å, which is more than enough for the smooth transportation of Li-ion between interstitial sites, and 2] the natural bonding between free Li-ion and anion [O⁻] network. The network anion has strong covalent bonding with all the three cations [Li⁺, Zn²⁺, Ge⁴⁺], particularly with Li⁺, which leads to the polarization of oxygen charge density away from interstitial ions and thus weakening the bonding with mobile Li-ion. The weak bonding between cations and oxygen anions further supports the fast movement of Lithium-ion. It is important here to note that due to the parallelogram shape of the bottlenecks, the lithium ions in Li₁₄Zn [GeO₄]₄ can diffuse only in two dimensions. ²² It has been observed that due to lower migration energy [0.21 eV–0.35 eV], the interstitial diffusion of Li-ions in LISICON microstructure is dominant over vacancy diffusion with migration energy of 0.56 eV. ²³ Even at such lower migration energy, the doped LISICON compounds, such as γ-Li_14.4 V_1.6Ge_2.4O₁₆, show relatively poor ionic conductivity [10⁻⁶ S cm⁻¹] as compared to that of liquid electrolytes [10⁻² S cm⁻¹] at room temperature. Also, LISICONs react with lithium metal and atmospheric gases [such as CO₂], which degrades their performance. Lower ionic conductivity and reactivity with ambient atmosphere hinder the use of LISICONs in ASSBs and other devices. ¹⁷

Therefore, to increase the ionic conductivity of LISICONs, Noriaki Kamaya et al. ²⁴ replaced oxygen with sulfur to develop this LISICON, Li₁₀GeP₂S₁₂, having ionic conductivity in the of order 10⁻² S cm⁻¹ at room temperature. This family is to be discussed in detail in the upcoming sulfide section.

In solid-state inorganic chemistry, Perovskite structure [ABO₃] with space group Pm $\bar{3}$ m [space group number 221], is encountered almost everywhere. This structure was reported by Latie et al. ²⁵ for the first time. It holds most of the metallic ions of the periodic table with a considerable number of different anions. The cubic unit cell of perovskite is shown in Fig. 4, ²⁶ where A-site ions [usually alkaline-earth or rare-earth elements] sits at the center of the cell, transition metal ions B sits at the corners, whereas the oxygen atoms sits at the face-centered positions of the corer shared octahedra. Here it is important to note that A sites are in 12-fold coordination and B sites are in 6-fold coordination oxygen anios. ²⁷

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The lithium-ion conductivity in perovskite structured Li_3xLa_2/3−xTiO₂ [LLTO] compounds was studied by Inaguma et al. ^28–30 Oguni et al. ³¹ and Kawai et al. ³² These authors have shown that in perovskites structures, the cation deficiency present at sites-A favors fast ionic conduction through the bottlenecks formed by four nearby octahedra, BO₆. The A and B sites of perovskites can hold different ions of different valence states. To optimize the best ionic conductivity, several substitutions have been done for sites A and B, ³³ but Inaguma et al. ²⁸ achieved the best results for Li_3xLa_2/3−xTiO₂ compounds. At room temperature for Li_0.34[1]La_0.51[1]TiO_2.94[2], they have obtained bulk ionic conductivity of 1 × 10⁻³ S cm⁻¹ and total ionic conductivity greater than 2 × 10⁻⁵ S cm⁻¹.

In the crystal structure of this compound, Li and La ions are in 12—fold coordination occupying the central sites A, and Ti-ions are in 6-fold coordination occupying the corner sites B. Whereas, bottleneck oxygen-ions create a potential barrier for Li-ions while they move from one A site to an adjacent A site. Previously we have discussed that the small or large size of bottleneck is responsible for the slow or fast movement of Li-ions in solid electrolytes respectively. Hence to obtain higher Li-ion conductivity many attempts have been made to increase the size of the bottleneck, either by changing the composition of elements or by substituting another element in the perovskite structure. Inaguma et al. ²⁹ substituted site A of LLTO by Sr of mole 5% and obtained bulk ionic conductivity 1.5 × 10⁻³ S cm⁻¹ at RT. This substitution increased the lattice constant which further increased the size of a bottleneck. The increased size of the bottleneck offered a convenient path for the movement of Li-ion and hence the Sr doped LLTO showed higher ionic conductivity than pure LLTO. However, doping of Sr beyond 10 mole % decreased the ionic conductivity due to the decreased concentration of Lithium ions.

In 2012, Zhao et al. ³⁴ presented a new class of 3D structured "lithium-rich anti-perovskites [LiRAP]" solid-state electrolyte. Unlike perovskites [A⁺B²⁺O⁻ ₃], to obtain Li-ion conductors, Zhao et al. did electronic inversion of conventional perovskites and replaced cations from anions and vice-versa to find anti-perovskite [A⁻B²⁻O⁺ ₃]. The XRD reveals typical [Pm $\bar{3}$ m] perovskite structure of anti-perovskites in which strongly electropositive monovalent cation [Li⁺] sits at the octahedral vertices, divalent anion sits at the octahedral center [site- B] whereas monovalent anion sits at the dodecahedral center [site- A]. In this structure, oxygen is the convenient choice for site-B and any halogen [F, Cl, Br, I] or a mixture of halogens occupies site-A. The reported ionic conductivity for Li₃OCl and Li₃OCl_0.5Br_0.5 is 0.85 × 10⁻³ S cm⁻¹ and 1.94 × 10⁻³ S cm⁻¹ respectively (at room temperature). Higher temperature causes a higher number of vacancies for site hopping of lithium ions in these compounds, leading to their higher ionic conductivity. Though excellent ionic conductivities have been obtained for Li-rich anti-perovskites Li_3−xOH_xCl, its lithium and proton mobilities, as a function of composition, are not completely characterized. James A. Dawson et al. ³⁵ were the first to study the mobility of lithium ions and protons in Li_3−xOH_xCl. With the help of ab initio molecular dynamics and ¹H, ²H, and ⁷Li solid-state NMR spectroscopy they have predicted an exothermic hydration enthalpy for Li₃OCl. It explains the complexities involved in the synthesis of moisture-free samples and the easiness with which the samples absorb the moisture.

This family of solid-state electrolytes is getting special attention due to the advantages such as higher ionic conductivity of about 10^–3 S cm⁻¹ at room temperature, excellent environmental stability, wide range electrochemical window [>6 V vs Li/Li⁺], and most stable interface against Li metal ${\rm{a}}$mong all the SSEs. ^36,37 The group of ortho-silicates with general formula A₃ ^IIB₂ ^III [SiO₄]₃ [A = Ca, Mg; B = Al, Cr, Fe], are considered as ideal garnets. It generally crystallizes in cubic structure [space group Ia $\bar{3}$ d] in which Si cations sits at a four-fold coordination site. Whereas A & B sits at eight & six-fold coordination sites respectively. Exponential growth in the ionic conductivity has been observed in the electrolytes of the group of this group with the increase in Li ratio. The higher concentration of Li leads to an organized increment in the population of octahedral sites, at the same time it introduces the vacancies at the tetrahedral sites. The increased octahedral population results in higher Li⁺-Li⁺ electrostatic repulsion. The electrostatic repulsion pushes the central octahedral lithium ions away to the adjacent tetrahedra. This inspection suggests relatively facile hoping of lithium ions in Li-stuffed garnet structure. ³⁸ Keeping it into sight, garnet-type electrolytes have been divided into four subgroups i.e. Li₃, Li₅, Li₆, and Li₇. Ca₃A_l2Si₃O₁₂ structure was the inspiration for the development of Li-containing garnets Li_xM₂M'₃O₁₂. ³⁷ Analogous to Ca₃A_l2Si₃O₁₂ structure, O'Callaghan et al. ³⁹ developed Li₃Ln₃Te₂O₁₂ [Ln = Y, Pr, Nd, Sm−Lu] compounds by replacing tetrahedral Si⁴⁺ by Li. These compounds, for example, Li₃Nd₃Te₂O₁₂, shows an ionic conductivity of [1 × 10⁻⁵ S cm⁻¹ at 600^oC] and activation energy [1.122 (15)eV]. ³⁹ The poor ionic conductivity of these compounds was attributed to the deficiency of Li content. Therefore, to obtain high ionic conductivity researchers started to synthesize garnets of higher Li content i.e. lithium stuffed garnets. ³⁷ In this effort, Thangadurai et al. ⁴⁰ developed the first Li stuffed garnet with the general formula Li₅La₃M₂O₁₂ [M = Nb, Ta] in 2005. The correct crystal structure of this compound has been quite controversial, especially because of the position of Li ions. Using neutron diffraction [ND], Cussan ⁴¹ proved that Li₅La₃M₂O₁₂ crystallizes in cubic structure [group symmetry Ia $\bar{3}$ d] with lattice constant a = 12.797 Å and 12.804 Å for Nb and Ta respectively. Also, tetrahedral and distorted octahedral sites hold the excess Lithium ions and La³⁺ & Nb⁵⁺ ions positioned at the eight and six coordination sites respectively. ^40,41 As compared with the ideal garnet composition, these compounds have an excess of 16 lithium atoms, and therefore these compounds, for example, Li₅La₃Nb₂O₁₂, exhibit higher bulk lithium-ion conductivity 8.0 × 10⁻⁶ S cm⁻¹ [At 22 °C] than Li₃ subtype garnets with activation energies of 0.43 at less than 300 °C. Even Greater Lithium-ion concentration (and hence higher Lithium-ion conductivity) can be achieved by partially substituting La³⁺ with divalent ions in Li₅La₃M₂O₁₂. This substitution leads Li₅La₃M₂O₁₂ to Li₆ phase compound with general formula Li₆ALa₂M₂O₁₂ [A = Mg, Ca, Sr, Ba, and M = Nb, Ta]. These compounds crystalize in the same cubic structure as their parent compound Li₅La₃M₂O₁₂. At 22 °C Li₆ALa₂M₂O₁₂ compounds, for example, Li₆BaLa₂Ta₂O₁₂ displayed ionic conductibility 4 × 10⁻⁵ S cm⁻¹, which is of course greater than previously, discussed lesser stuffed garnet subtypes Li₃ and Li₅, and activation energy 0.40 eV. ^40,42 In 2007, Murugan et al. ⁴³ synthesized the fourth subtype of garnet structured material "Li₇La₃Zr₂O₁₂" by replacing M in the second subtype garnet structured compound Li₅La₃M₂O₁₂ from Zr. This new compound has a Li₇ phase that shows the highest ionic conductivity of 3 × 10⁻⁴ S cm⁻¹ and lowest activation energy of 0.3 eV among all the previously discussed garnet subtypes namely Li₃, Li_5, and Li₆. For charge balancing, this phase was stuffed with excessive Li-ions. Two structural types are associated with Li₇La₃Zr₂O₁₂; a cubic garnet, having the same standard pattern as that of previously discussed garnet phase Li₅La₃Nb₂O₁₂ excepting subtle changes in the symmetry and a tetragonal structure with space group I4₁ /acd [no. 142] and lattice parameters a = 13.134[4] Å and c = 12.663[8] Å. ^43,44

In 2010, Ramzy et al. ⁴⁵ investigated the lithium ionic conductivity of the solid solutions Li₇La₃Zr₂O₁₂ and Li₆La₂BaM₂O₁₂ [M = Nb, Ta]. The solid solutions were taken in ratio 1:1 and using solid-state reaction the compound Li_6.5La_2.5BaZrTaO₁₂ was prepared which showed an ionic conductivity 6 × 10⁻³ S cm⁻¹ at 100 °C. Despite around two decades of efforts to obtain the Li-ion conductivity comparable to that of organic electrolytes 1 × 10⁻² S cm⁻¹ at RT the current highest ionic conductivity has been achieved by Qin et al. ⁴⁶ In 2018 they have synthesized the highly textured Ga₂O₃-substituted Li₇La₃Zr₂O₁₂ compound Li_6.55Ga_0.15La₃Zr₂O₁₂ that showed the ionic conductivity of 2.06 × 10⁻³ S cm⁻¹ at RT, which is still far away from that of organic electrolytes at RT. To fight the high resistance and low capacity of ASSBs, Gregory T. Hitz et al. ⁴⁷ fabricated a unique Li₇La₃Zr₂O₁₂ [LLZO] doped ceramic microstructure. With the help of a scalable roll-to-roll manufacturing technique, they prepared a porous-dense-porous tri-layer structure. The Lithium symmetric cells based on this structure were cycled at RT and achieved dramatically lower area-specific resistances [∼7 Ω-cm²], and dramatically higher current densities [10 mA cm⁻²]. Dawei Wang et al. ⁴⁸ used the co-sintering method to construct Solid-state lithium batteries [SSLBs]. To rope, the cathode materials LiNi_0.6Mn_0.2Co_0.2O₂ [NMC] and solid-state electrolytes Li_6.4La₃Zr_1.4Ta_0.6O₁₂ they used Li₃BO₃ as a sintering agent. They obtained SSLBs of higher capacity [106 mAh g⁻¹] with small NMC primary particles as compared to that of prepared with large NMC secondary particles. Better interfacial properties and shorter Li-ion diffusion paths within the small-NMC were attributed to the higher capacity of SSLBs. With 0 to 9 V electrochemical stability window, Gregory T. Hitz et al. ⁴⁹ developed Ca doped garnet structured Li_7.1La₃Zr_1.95Ca_0.05O₁₂ solid state electrolyte. At room temperature it has shown the high ion conductivity and reduced activation energy of 5.2 × 10⁻⁴ S cm⁻¹ and 0.27 eV, respectively. In order to enhance ionic conductivity and suppress the lithium metal dendrite formation, Shufeng Song et al. ⁵⁰ developed multi-substituted garnet-type solid electrolyte Li_6.4Ga_0.1La₃Zr_1.55Ba_0.05Ta_0.4O₁₂. Apart from the good ionic conductivity (1.02 × 10⁻³ S cm⁻¹ at RT), the symmetric cell prepared with it (Li/Li_6.4Ga_0.1La₃Zr_1.55Ba_0.05Ta_0.4O₁₂/Li) exhibits excellent cycling stability.

As discussed in the previous sections, an ideal solid-state electrolyte for all-solid-state lithium-ion secondary batteries should have high ionic conductivity, excellent chemical stability against electrode materials, moisture, and atmospheric gases. Low cost, ease of availability of the starting materials, and thermal stability are other prerequisites of solid-state electrolytes. Li₇La₃Zr₂O₁₂ surprisingly fulfill all the above precondition which makes it one of the most ideal electrolytes for ASSBs. ⁴³

The structure of this generation of electrolytes differs from that of other electrolytes discussed above. These electrolytes have disordered or network structures as shown in Fig. 5, ⁵¹ and therefore are referred to as amorphous/glass type solid-state electrolytes.

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Lithium Phosphate Oxy-Nitride [LiPON] is an example of an Amorphous/glass type solid-state electrolyte that was discovered by Bates and co-workers in 1992. Usually, it is prepared in a nitrogen atmosphere by sputtering a crystalline Li₃PO₄ target. The nitrogen atmosphere helps in the formation of phosphorus–nitrogen bonds, which enhances the ionic conductivity of LiPON [Li_3.3PO_3.9N_0.17] to 2 × 10⁻⁶ S cm⁻¹ at room temperature. ⁵² Since the discovery, many efforts have been done to improve the ionic conductivity of LiPON by either increasing the amount of lithium-ion in the target mixture up to a certain amount ⁵³ or by other methods such as by adjusting the ratio of the flow rate of N₂ and Ar, ⁵⁴ or by increasing the N₂ content, ⁵⁵ or by changing the methods of film preparation, ⁵⁶ but unfortunately, the ionic conductivity of Li ions couldn't be increased beyond the order of 10⁻⁵ cm S⁻¹, at room temperature. The electrolytes of this family are considered promising for thin-film ASSBs and inappropriate for bulk ASSBs. ^53,57,58

LiSON ⁵⁹ and LiPOS [6LiI–4Li₃PO₄–P₂S₅] ⁶⁰ are other glassy systems. Similar to LiPOS systems, the LiSON composition Li_0.29S_0.28O_0.35N_0.09 shows an ionic conductivity of 2 × 10⁻⁵ S cm⁻¹ at room temperature. Other systems like LiBSO [0.3 liBO₂–0.7 li₂SO₄] ⁶¹ or LiSiPON [Li_2.9Si_0.45PO_1.6N_1.3] also belong to this family of solid-state electrolytes, ⁶² having poor ionic conductivity and cyclability.

With the Li₂S–SiS₂ system, the research into sulfide-type solid-state electrolytes started in 1986. ⁶³ Kennedy et al. ⁶⁴ were the first to report Li⁺-ion conduction in a melt-quenched glass of Li₂S-SiS₂ up to the order of 10⁻⁴ S cm⁻¹ at room temperature. Since then, to improve the ionic conductivity and electrochemical stability, this family of solid-state electrolytes has been studied extensively. However, doping is proved to be the most efficient technique to improve ionic conductivity. Kondo et al. ⁶⁵ used the liquid nitrogen quenching method to dope Li₂S-SiS₂ glass by Li₃PO₄ and obtained improved ionic conductivity and electrochemical stability than pure Li₂S–SiS₂ system at room temperature. Similarly, Fuminori Mizuno et al. ⁶⁶ doped Li₂S–SiS₂ glasses with a small amount of Li₄SiO₄ and obtained the Li⁺ ion conductivity around 10⁻³ S cm⁻¹ and higher stability as compared to that of pure Li₂S–SiS₂ glasses. Hayashi et al. ⁶⁷ prepared oxy-sulfide amorphous materials in the systems 95[0.6Li₂S-0.4SiS₂].5Li_xMO_y [M = Si, P, and Ge] by melt-quenching and mechanical milling techniques. At room temperature, these materials exhibited ionic conductivities over the order of 10⁻³ S cm⁻¹ and have also shown a wide potential window of 10V. The laboratory-scale solid-state batteries using oxy-sulfide amorphous solid electrolytes are one of the most suitable solid-state lithium-ion batteries exhibiting excellent cycling performance up to 100 times. ⁶⁷

This family of electrolytes seems to be promising due to higher ionic conductivity [up to the order of 10⁻³ S cm⁻¹], low activation energy [18 kJ mol⁻¹], and higher electrochemical stability with the electrodes. ^68–71 These are prepared through various methods such as melt quenching, ball milling, and solid-state reactions. In past decades melt quenching method has been a favorable option for the preparation of these sulfides. In this method, due to the high vapor pressure of P₂S₅, the melting is performed in sealed quartz tube. ^68,72,73 Melt-quench method was used by Zhang et al. ⁷² to prepare 0.33[[1-y] B₂S₃-yP₂S₅]−0.67Li₂S [0≤y≤0.3, 0.9≤y≤1.0] glasses. For sample 0.23B₂S₅−0.10P₂S₅− 0.67Li₂S [y = 0.3], they have reported an ionic conductivity of 1.41 × 10⁻⁴ S cm⁻¹ at room temperature. Using mechanical milling and solid-state reaction method, Mizuno et al. ⁷⁴ have developed the ceramic, glass, and glass-ceramic 70Li₂S-30P₂S₅ solid-state electrolytes. The glass and glass-ceramic samples, prepared by the mechanical milling process showed up good ionic conductivity [5.4 × 10⁻⁵ S cm⁻¹ and 3.2 × 10⁻³ S cm⁻¹ respectively] and low activation energy [38 kJ mol⁻¹ and 18 kJ mol⁻¹ respectively] at room temperature. Whereas the ceramic sample developed by solid-state reaction method has shown poor ionic conductivity of 2.6 × 10⁻⁸ S cm⁻¹ and large activation energy 55 kJ mol⁻¹ at room temperature. Hence, to prepare Li₂S-P₂S₅ solid electrolytes of low activation energy and higher ionic conductivity, the solid-state reaction isn't an appropriate method. ⁷⁴ The higher conductivity of the samples obtained through the ball milling process is attributed to the production of ultra-fine amorphous materials. At the same time materials obtained via this method form close contact with electrodes, leading to the superior electrochemical performance of ASSBs. ⁷⁵ Here it is important to note that due to precipitation of low conductive crystals, in general, the ionic conductivity of glasses decreases as their crystallinity increases. However, proper composition and heat treatment of Li₂S-P₂S₅ glasses lead to the precipitation of crystalline super-ionic metastable phase which enhances its ionic conductivity. ⁷⁶ The metastable phases are only precipitated from glass and could not be obtained from solid-state reactions. Too high temperature leads to less conductive thermodynamically stable phases [Li₄P₂S₆ for 70Li₂S-30P₂S₅ and Li_3.55P_0.89S₄ for 80Li₂S- 20P₂S₅], however, to obtain and maintain highly ionic conductive metastable phases the temperature of the glasses is raised just above the first crystallization temperature. ⁷⁷

Using synchrotron X-ray powder diffraction, Yamane et al. ⁷⁷ determined a new and highly conductive phase - Li₇P₃S₁₁. This phase consists of the two most suitable structural units; Pyro-thiophosphate [P₂S₇ ⁴⁻] and ortho-thiophosphate [PS₄ ³⁻], which are responsible for the fast lithium-ion conduction in Li₂S-P₂S₅ glass-ceramics. To obtain Li₂S-P₂S₅ glass-ceramic with metastable phase Li₇P₃S₁₁, the heat treatment conditions were optimized by Seino et al. ⁷⁸ Due to negligibly present grain boundaries in the microstructure [and hence better contact between grains], the Li₂S-P₂S₅ glass-ceramic showed up higher ionic conductivity of magnitude 1.7 × 10⁻² S cm⁻¹ at 25 °C.

Although, Li₂S-P₂S₅ glass/glass-ceramics have shown the ionic conductivity reaching that of organic electrolytes, still challenging to use them in an open atmosphere because the water molecule present in the air hydrolyses them and generate harmful H₂S gas. The generation of harmful H₂S gas is accompanied by structural changes, though the amount of H₂S production and the extension of structural change has been observed different for several Li₂S-P₂S₅ electrolytes. In the case of 67Li₂S33P₂S₅ glass and Li₂S crystal, significant H₂S was generated and surprisingly great structural change was reported in the units of S²⁻ or P₂S₇ ⁴⁻ ions. On the other hand, the structural unit of PS₄ ³⁻ ion present in 75Li₂S25P₂S₅ glass and glass-ceramic did not show any significant structural change and generates the least amount of H₂S gas on exposing it to the open atmosphere. Therefore, in terms of chemical stability and ionic conductivity, the 75Li₂S25P₂S₅ glass and glass-ceramic are favorable solid-state electrolytes for ASSBs. ⁷⁹

In an attempt to achieve higher ionic conductivities a series of sulfide crystals, called thio-LISICON, have been obtained by Kanno and his co-workers by just replacing O²⁻ by S²⁻ in LISICONs. Li_4−xM_1−yM'_y S₄ [where the M and M' are Ge, Si or P, Al, Ga, respectively] is the general composition of this family. Similar to the oxide counterparts, the thio-LISICONs also have γ–Li₃PO₄ framework structure but these have shown higher ionic conductivities (̴ 10⁻³ S cm⁻¹ at 25 °C] due to weaker force of attraction between S²⁻ and Li⁺ than that of between O²⁻ and Li⁺ in LISICONs. ^80–83 Among six new compounds [Li₂GeS₃, Li₄GeS₄, Li₂ZnGeS₄, Li_4–2xZn_xGeS₄, Li₅GaS₄ and Li_4+x+δ [Ge_1−δ'−xGa_x]S₄] prepared by Kanno et al. ⁸⁰ Li_4+x+δ [Ge_1−δ'−xGa_x]S₄] has shown highest ionic conductivity of 6.5 × 10⁻⁵ S cm⁻¹ [at 25 °C for x = 0.25] among all. Also, it has a wide electrochemical window of 5 V vs Li/Li⁺.

Vacancy and interstitial doping are the two doping methods to improve the ionic conductivities of thio-LISICONs based on Li₄SiS₄ and Li₄GeS₄. A dramatic enhancement in conductivities is observed when these are doped to create either Li⁺ vacancies or Li⁺ interstitials. New thio-LISICON solid electrolyte, Li_4−xGe_1−xP_xS₄, was prepared by partial vacancy doping of P⁵⁺ at the site Ge⁴⁺ in the system Li₄GeS₄. Maximum ionic conductivity [2.17 × 10⁻³ S cm⁻¹] and minimum activation energy [20 kJ mol⁻¹] were reported at 25 °C for x = 0.75. ⁸² In 2011, Kamaya and his co-workers synthesized a new compound of thio-LISICON family. At 27 °C, that new compound Li₁₀GeP₂S₁₂ showed an ionic conductivity of 1.2 × 10⁻² S cm⁻¹. However, this ratio could fit the earlier discussed thio-LISICON, Li_4−xGe_{1− x}P_xS₄ when x = 2/3, but it is structurally different from the previously synthesized thio-LISICON compound Li_4−xGe_1−xP_xS₄ for x = 0.75 i.e. Li_3.25Ge_0.25P_0.75S₄. ²⁴ Figure 6 ²⁴ represents the three-dimensional crystal structure of Li₁₀GeP₂S₁₂. It consists of [Ge_0.5P_0.5] S₄ tetrahedrons, PS₄ tetrahedrons, LiS₄ tetrahedrons, and LiS₆ octahedrons. The neutron diffraction studies of Li₁₀GeP₂S₁₂ have revealed that one-dimensional conduction pathways are followed by lithium ions within the LiS₄ tetrahedrons. In the c direction, the predicted ionic conductivity was 4 × 10⁻² S cm⁻¹, whereas that of in a- b was predicted 9 × 10⁻⁴ S cm⁻¹ at 27 °C. Since the Li-ions can crossover between 1-dimensional channels in c direction therefore for overall diffusion, the diffusion in a-b plane is also important. So we can say that the good ionic conductivity of Li₁₀GeP₂S₁₂ is the consequence of 3D diffusion pathways along c axis as well as in a-b plane. At room temperature, the ionic conductivity of Li₁₀GeP₂S₁₂ approaches that of organic electrolytes i.e. ̴ 10⁻² S cm⁻¹. Even at low temperatures, for example at − 30 °C and − 45 °C, Li₁₀GeP₂S₁₂ has shown an ionic conductivity of 1 × 10⁻³ S cm⁻¹ and 4 × 10⁻⁴ S cm⁻¹ respectively. Also, Li₁₀GeP₂S₁₂ has wide electrochemical of 5 V vs Li/Li⁺. ^24,84 However, due to limited deposits [and therefore higher cost of germanium], the application of LGPS could be practically challenging. In an attempt to fight the higher cost of germanium Bron et al. ⁸⁵ successfully synthesized a novel material just by replacing germanium from inexpensive tin [Sn] and obtained Li₁₀SnP₂S₁₂ having an ionic conductivity of 7 × 10⁻³ S cm⁻¹ at 27 °C. The replacement of germanium by tin may reduce the cost by a factor of 3.

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Later, to enhance the ionic conductivity of Li₁₀GeP₂S₁₂, it was doped with oxygen and chlorine. Chlorine doped Li₁₀GeP₂S₁₂ i.e. Li_9.54Si_1.74P_1.44S_11.7Cl_0.3 has shown the highest ionic conductivity of 2.5 × 10⁻² S cm⁻¹ at room temperature among all the reported Li-ion electrolytes so far. ^86,87 Though, ASSBs are considered promising in terms of energy density for next-generation energy storage systems. Nevertheless, their insufficient ionic and electronic percolation within the composite cathode limits the performance. ⁸⁸

Lithium is considered the most suitable metal for the anode of high-energy secondary batteries. But the dendrite formation hinders the use of Li metal as anodes with traditional liquid or polymer electrolytes. However, due to high mechanical strength and high Li-ion transference number, the solid-state electrolytes are assumed to be suitable for the prevention of Li dendrite growth. Due to higher ionic conductivity and excellent compatibility with lithium metal anodes, Li₇La₃Zr₂O₁₂ [LLZO] and Li₂S–P₂S₅ is considered the most promising solid-state electrolytes. Nevertheless, recent reports have shown the formation of lithium dendrites in LLZOs and Li₂S–P₂S₅ solid-state electrolytes. ⁸⁹ With the help of Li NMR chemical shift imaging and electron microscopy, Lauren E. Marbella et al. ⁹⁰ monitored the primary stages of dendrites growth in the garnet-type solid electrolyte, Li_6.5La₃Zr_1.5Ta_0.5O₁₂. The inorganic solid-state electrolytes with their composition, ionic conductivity, and corresponding temperature have been summarised in the Table I below.