DFT Modelling of Li₆SiO₄Cl₂ Electrolyte Material for Li-Ion Batteries

Abstract

There is significant interest in finding a promising lithium-containing oxide that can act as a solid electrolyte in a rechargeable lithium-ion battery. Li₆SiO₄Cl₂ is a candidate electrolyte material which was recently characterized using both experimental and computational techniques. In this study, density functional theory simulation was used to examine the intrinsic defects, solution of promising isovalent and aliovalent dopants, possible reaction routes for the formation of Li₆SiO₄Cl_2, and the feasibility of incorporating additional Li in this material. The results revealed that the O–Cl anti-site cluster was the lowest energy defect in this material. The LiCl Schottky was the second lowest energy defect process, and the Li Frenkel was higher—only by 0.06 eV—than the LiCl Schottky. The candidate dopants on the Li, Si and Cl were Na, Ge and F, respectively. Substituting Al on the Si site was an efficient way of increasing the amount of Li in this material. Incorporation of extra Li (up to three) was considered and this process was endothermic. Different chemical reaction routes were constructed and their reaction energies were calculated to predict the feasibility of the formation of Li₆SiO₄Cl₂. The formation of Li₆SiO₄Cl₂ from constituent elements (Li, Si O₂ and Cl₂) is thermodynamically feasible.

1. Introduction

The generation of efficient energy storage appliances such as batteries, capacitors and fuel cells is a key task, as there is a global energy demand in the industrial and technological market due to the growing human population. Furthermore, these devices are alternative sources of energy to replace fossil fuels, which will possibly run out in a couple of decades.

Lithium-ion batteries (LIBs) [1,2,3,4,5] have drawn significant research interest owing to their applications in many everyday electronic appliances such as laptops, mobile phones and electrical vehicles. LIBs have many positive characteristics such as excellent efficiency, cycle performance, high power, and energy densities, although they require protection during intercalation (charge/discharge) and are expensive [6,7,8]. The efficacy of LIBs is dependent on the performance of the materials used for the preparation of anodes, cathodes and electrolytes. Furthermore, materials that are cheap, non-hazardous, safe and of high abundance, are of great interest for designing a promising Li-ion battery.

Solid electrolytes are key components in the construction of LIBs, and there are safety risks of using liquid electrolytes such as leakage and flammability [9,10,11]. The properties of the electrolyte have a significant impact on the performance of a battery. A variety of solid electrolyte materials such as Li_3xLa_(2/3)−xTiO₃ [12]_, Li₃OCl [13], LiZr₂(PO₄)₃ [14], Li₇La₃Zr₂O₁₂ [15], Li₁₀GeP₂S₁₂ [16], and LiPON [17,18] have been assessed both experimentally and theoretically. There are necessary requirements (low electronic conductivity, high Li-ion diffusion, feasible electrochemical stability, and good interfacial contact with electrodes) that should be satisfied before commercial application. The garnet-type Li₇La₃Zr₂O₁₂ has received great attention recently due to its high ionic conductivity. The tetragonal phase of Li₇La₃Zr₂O₁₂ exhibited higher Li-ion conductivity at low temperatures than that measured in its cubic phase [15]. The stability of the cubic phase and its ionic conductivity at high temperatures were enhanced by doping of Al, Ga, Ta and Nb [19]. The activation energy of Li-ion diffusion was calculated experimentally by Shao et al. [15] and its low value of 0.25 eV was linked to high ionic conductivity.

Li₆SiO₄Cl₂ is a candidate solid electrolyte material that belongs to the cubic lithium argyrodite family [20]. This material was recently synthesized and the ionic conductivity of this material was reported to be low at room temperature [20]. The structural rigidity of this material is provided by the strong [SiO₄]⁴⁻ units. Furthermore, six Li can theoretically be extracted from each per formula unit, although the exact number is not known experimentally. As this material was recently established, there are no further experimental or theoretical studies available in the literature. In particular, theoretical studies of the intrinsic defect chemistry, solution of dopants, synthetic routes and Li-incorporation can provide further valuable information about this material.

In this study, materials modelling technique based on the density functional theory (DFT) was used to examine defect chemistry, solution of dopants, different formation reaction routes and Li-incorporation. This simulation approach provided electronic structures of promising doped and Li-incorporated configurations. In previous simulation studies, the same technique was successfully applied to various electrolyte and electrode materials [18,21,22].

2. Computational Methods

DFT simulations, together with spin polarization, were employed using the VASP (Vienna ab initio Simulation Package) code [23]. In all cases, projected augmented wave (PAW) pseudopotentials [24], and a plane wave with a basis set consisting of a 500 eV cut-off, were used. Bulk Li₆SiO₄Cl₂ was modelled using an 8 × 8 × 8 Monkhorst–Pack [25] k-point mesh. All defect calculations were performed using a 1 × 2 × 2 super cell consisting of 208 atoms (Cl: 32, Si: 16, O: 64 and Li: 96). The outer shell configurations for Li, Si, Cl and O, were 2s¹, 3s² 3p², 3s² 3p⁵ and 2s² 2p⁴, respectively. For defect structures, a 4 × 4 × 4 Monkhorst–Pack k-point mesh was used. The exchange-correlation term was defined by the aid of the generalized gradient approximation (GGA) scheme as formulated by Perdew, Burke and Ernzerhof (PBE) [26]. We used the conjugate gradient algorithm [27] to optimize the structures completely with a force tolerance which was less than 0.001 eV/Å. The charges on the atoms of the optimized structures were calculated using the Bader charge analysis [28]. In this analysis, zero flux surfaces were used to divide molecules into atoms. Density functional theory calculations provided structural information together with electronic structures.

3. Results and Discussion

3.1. Bulk Structure of Li₆SiO₄Cl₂

A recent combined experimental and theoretical study showed that although Li₆SiO₄Cl₂ can exist in different phases, the most energetically stable phase is orthorhombic (space group Pna2₁) [20]. Its lattice constants are a = 10.5204 Å, b = 6.0756 Å, c = 9.9530 Å and α = β = γ = 90.0° [20]. In this structure, Li, Cl and Si form tetrahedral (LiO₂Cl₂), octahedral (ClLi₆) and tetrahedral (SiO₄) environments, respectively (see Figure 1). The current simulation study considered the most stable phase reported in the previous study [20], although the formation of other configurations of Li₆SiO₄Cl₂ beyond the orthorhombic phase could be examined. Full geometry optimisation of this structure was carried out to compare calculated and experimentally observed structures.

Table 1. Calculated and experimental parameters of Li₆SiO₄Cl₂.

Parameter	Calculated	Experiment [20]	|∆| (%)
a (Å)	10.479	10.524	0.43
b (Å)	6.036	6.076	0.66
c (Å)	9.918	9.953	0.35
α = β = γ (°)	90.0	90.0	0.00
V (Å3)	627.34	636.17	0.16

shows the calculated and experimental lattice parameters and volume. An excellent agreement was observed (see Table 1), showing the efficacy of the pseudo-potentials and basis sets employed in this study. The percentage error calculated for each parameter was less than 1. The density of state (DOS) plot of bulk Li₆SiO₄Cl₂ is shown in Figure 1b. It is an insulator with a band gap of 6.00 eV, showing the desirability of this material as an electrolyte material in Li-ion batteries. The electronic structure of this material was absent in the literature to compare the band gap value calculated in this study.

A Bader charge analysis was carried out to calculate the charges on the atoms in the bulk. The results tabulated in

Table 2. Calculated Bader charges on each atom present in the bulk Li₆SiO₄Cl₂.

Atom	Bader Charge
Li	+1.00
Si	+4.00
O	−1.99, −2.01
Cl	−0.99

clearly show that the crystal structure was ionic, with all ions consisting of full charges (Li⁺, Si⁴⁺, O²⁻ and Cl⁻).

3.2. Intrinsic Defect Properties

Intrinsic defect processes are important as they influence the properties of materials. Here, the formation of Frenkel, Schottky and anti-site defect processes are examined. These defect processes require point defects together with their energies, as explained in the following equations consisting of Kröger–Vink notations [29]. While the Schottky defect produces many vacancies, in Frenkel defects, ions occupy interstitial positions. Both Schottky and Frenkel defect concentrations are temperature-dependent, and are generally expected to increase with increasing temperature [30]. Ionic conductivity of a material is controlled by Frenkel defects. In an experimental study, the mechanism for the high conductivity of Ag⁺ ions in all phases of AgI was explained using Frenkel defects [31]. The electronic conductivity enhancement was linked to the formation of Schottky defects in LaFeO₃.

First, point defect energies were computed and associated to calculate Frenkel, Schottky and anti-site defect formation energies. The calculation formula for each defect process was reported in our previous simulation study [32] and is provided in the electronic supplmentary information (ESI). Ths total energies of point defects are also provided in the ESI (see Table S1).

$$\mathrm{Li}\mathrm{Frenkel}:{\mathrm{Li}}_{\mathrm{Li}}^{\mathrm{X}}\to V_{\mathrm{Li}}^{\prime }+{\mathrm{Li}}_{\mathrm{i}}^{•}$$

(1)

$$\mathrm{Si}\mathrm{Frenkel}:{\mathrm{Si}}_{\mathrm{Si}}^{\mathrm{X}}\to V_{\mathrm{Si}}^{\prime \prime \prime \prime }+{\mathrm{Si}}_{\mathrm{i}}^{••••}$$

(2)

$$\mathrm{Cl}\mathrm{Frenkel}:{\mathrm{Cl}}_{\mathrm{Cl}}^{\mathrm{X}}\to V_{\mathrm{Cl}}^{•}+{\mathrm{Cl}}_{\mathrm{i}}^{\prime }$$

(3)

$$\mathrm{O}\mathrm{Frenkel}:{\mathrm{O}}_{\mathrm{O}}^{\mathrm{X}}\to V_{\mathrm{O}}^{••}+{\mathrm{O}}_{\mathrm{i}}^{\prime \prime }$$

(4)

$$\mathrm{Schottky}:6{\mathrm{Li}}_{\mathrm{Li}}^{\mathrm{X}}+{\mathrm{Si}}_{\mathrm{Si}}^{\mathrm{X}}+4{\mathrm{O}}_{\mathrm{O}}^{\mathrm{X}}+2{\mathrm{Cl}}_{\mathrm{Cl}}^{\mathrm{X}}\to 6V_{\mathrm{Li}}^{\prime }+V_{\mathrm{Si}}^{\prime \prime \prime \prime }+4V_{\mathrm{O}}^{••}+2V_{\mathrm{Cl}}^{•}+{\mathrm{Li}}_6{\mathrm{SiO}}_4{\mathrm{Cl}}_2$$

(5)

$${\mathrm{Li}}_2\mathrm{O}\mathrm{Schottky}:2{\mathrm{Li}}_{\mathrm{Li}}^{\mathrm{X}}+{\mathrm{O}}_{\mathrm{O}}^{\mathrm{X}}\to 2V_{\mathrm{Li}}^{\prime }+V_{\mathrm{O}}^{••}+{\mathrm{Li}}_2\mathrm{O}$$

(6)

$$\mathrm{LiCl}\mathrm{Schottky}:{\mathrm{Li}}_{\mathrm{Li}}^{\mathrm{X}}+{\mathrm{Cl}}_{\mathrm{Cl}}^{\mathrm{X}}\to V_{\mathrm{Li}}^{\prime }+V_{\mathrm{Cl}}^{•}+\mathrm{LiCl}$$

(7)

$${\mathrm{SiO}}_2\mathrm{Schottky}:{\mathrm{Si}}_{\mathrm{Si}}^{\mathrm{X}}+2{\mathrm{O}}_{\mathrm{O}}^{\mathrm{X}}\to V_{\mathrm{Si}}^{\prime \prime \prime \prime }+2V_{\mathrm{O}}^{••}+{\mathrm{SiO}}_2$$

(8)

$$\mathrm{O}/\mathrm{Cl}\mathrm{anti}-\mathrm{site}\left(\mathrm{isolated}\right):{\mathrm{O}}_{\mathrm{O}}^{\mathrm{X}}+{\mathrm{Cl}}_{\mathrm{Cl}}^{\mathrm{X}}\to {\mathrm{O}}_{\mathrm{Cl}}^{\prime }+{\mathrm{Cl}}_{\mathrm{O}}^{•}$$

(9)

$$\mathrm{O}/\mathrm{Cl}\mathrm{anti}-\mathrm{site}\left(\mathrm{cluster}\right):{\mathrm{O}}_{\mathrm{O}}^{\mathrm{X}}+{\mathrm{Cl}}_{\mathrm{Cl}}^{\mathrm{X}}\to {\{{\mathrm{O}}_{\mathrm{Cl}}^{\prime }:{\mathrm{Cl}}_{\mathrm{O}}^{•}\}}^X$$

(10)

Defect formation energies calculated for all defects are listed in

Table 3. Calculated intrinsic defect energies for different defect processes in Li₆SiO₄Cl₂.

Defect Process	Reaction Equation	Reaction Energy/Defect (eV)
Li Frenkel	(1)	2.94
Si Frenkel	(2)	10.39
O Frenkel	(3)	5.96
Cl Frenkel	(4)	3.07
Schottky	(5)	3.72
Li2O Schottky	(6)	3.22
LiCl Schottky	(7)	2.88
SiO2 Schottky	(8)	6.52
O/Cl anti-site (isolated)	(9)	3.71
O/Cl anti-site (cluster)	(10)	0.07

. The O–Cl anti-site defect cluster was the lowest energy defect process, inferring a small amount of anion mixing was possible in this material (0.07 eV/defect). The diffusion property of this material could be affected by this defect. Individual defects (${\mathrm{O}}_{\mathrm{Cl}}^{\prime }{\mathrm{and}\mathrm{Cl}}_{\mathrm{O}}^{•}$) were considered separately in the isolated form of the anti-site defect. These two defects were simulated simultaneously in the cluster form. The formation of an anti-site defect cluster is exothermic (−3.64 eV), meaning that once the isolated defects are formed, there is a tendency to form cluster. Anti-site defect has been noticed in the experimental and theoretical characterization of many oxide materials [33,34,35,36]. We did not consider the Li-Si cation anti-site defect because the formation energy of this defect is expected to be highly positive according to our previous simulation study [37]. This is due to the large charge mismatch between Li⁺ and Si⁴⁺ ions. A cation anti-site defect containing Li can influence the Li diffusion properties. In an experimental study, it was shown that Li-ion migration pathways were affected by the anti-site defect consisting of Li and Fe in LiFePO₄ [38]. Kempaiah et al. [39] observed the presence of Li atoms on the Ni site and vice versa in LiNiPO₄ and discussed the electrochemical performance in the existence of an anti-site defect. The presence of an Li-Fe anti-site defect in the as-prepared Li₂FeSiO₄ was reported by Milović et al. [40] in their experimental and theoretical characterisation. This anti-site defect influenced the geometrical and electronic properties of Li₂FeSiO₄. The same anti-site defect was observed in Li₂FeSiO₄ only when this material was under cycling [41]. This clearly indicates that this defect is dependent on different synthetic conditions. Both Li-ion migration pathways and activation energies calculated for the cycled structure were different from those calculated for the as-prepared structure [41].

The next most favorable defect was the LiCl Schottky (2.88 eV/defect). This indicated that loss of LiCl is practicable at high temperatures only. The Li Frenkel defect energy was 2.94 eV/defect, meaning that the concentration of this defect was not very high. In previous simulation studies, the Li Frenkel was found to be the most dominant defect process [42,43]. The Cl Frenkel was more favorable than the O Frenkel. This indicated the concentration of Cl vacancies would be higher than that of O vacancies in this material. Other Frenkel and Schottky defect energies are highly endothermic, suggesting that they will not present at room temperature. In particular, the formation of Si Frenkel was hugely positive (10.39 eV/defect).

3.3. Solution of Dopants

A variety of metal oxides were examined to improve their properties via doping [44,45]. Yttria-doped zirconia is a promising ceramic electrolyte material with higher oxygen ion conductivity than pure zirconia [46]. The enhancement in the ionic conductivity is related to the presence of point defects controlling the transport properties at grainboundaries [46]. In a combined experimental and thoeretical study [47], it was demonstrated that substituting Ru on the Fe site in LiFePO₄ improved the Li-ion movement via reduction of the Li-Li separation in the lattice and electrical conductivity. Substitutional doping on the Li, Si and Cl sites can tune the properties of Li₆SiO₄Cl₂. Here, we considered both isovalent and aliovalent dopants. Necessary charge compensating defects (vacancies and interstitials) were introduced together with approriate lattice energies. The favourable dopants can be of interest for future experimental studies.

First, alkali metal dopants (R = Na, K, Rb and Cs) were doped on the Li site. Solution energies were calculated by using the following equation.

$${\mathrm{R}}_2\mathrm{O}+2{\mathrm{Li}}_{\mathrm{Li}}^{\mathrm{X}}\to 2{\mathrm{R}}_{\mathrm{Li}}^{\mathrm{X}}+{\mathrm{Li}}_2\mathrm{O}$$

(11)

Figure 2 shows the solution energies for the monovalent dopants. The most promising dopant was found to be the Na⁺. Exothermic solution energy (−0.74 eV) was calculated for this dopant. This was partly due to the smaller ionic radius of Na⁺ than that of the other dopant ions (see Figure 2). The next most favorable dopant was the K⁺. However, doping of this cation was endothermic. Once the ionic radius increased, solution energy also increased. The most inappropriate dopant was the Cs⁺, and this cation could not be doped at room temperature. The possible doped structural formula is Li_6−xNa_xSiO₄Cl₂ (x = 0.0–1.0). The doping of Na on the Li site was considered in an experimental study by Xue et al. [48] to reduce the concentration of Li in a mixed oxide Li_1.2Mn_0.54Ni_0.13Co_0.13O₂ as this overlithated material exhibited poor stability during cycling, and attenuation of voltage, although high capacity was achieved. There was an enhancement in the electrochemical property, cycling performance and reduction in voltage attenuation upon Na doping. As Li₆SiO₄Cl₂ is also an overlithiated material, doping of Na can be an effective method to improve the electrochemical properties of this material. In another experimental study by Shen et al. [49], the doping of Na in a “layered” LiNi_0.6Co_0.05Mn_0.35O₂ inhibited phase transition and cation mixing (Li/Ni), and improved the performance of Li storage.

The DFT-optimized structure of Na-doped Li₆SiO₄Cl₂ is shown in Figure 3a. There was an elongation in the bond distances (Na–O and Na–Cl) upon doping. This was due to the smaller ionic radius of Li⁺ than Na⁺. The electronic structure had a very small alteration upon doping, and the doped configuration was still an insulator. The p-states of Na lay closer to the Fermi level at 1.05 eV.

Next, trivalent dopants were considered on the Si site, as they are most interesting for practical application. This doping process introduced additional Li interstitials as charge compensating defects required for the enhancement in capacity, as defined by Equation (12).

$${\mathrm{R}}_2{\mathrm{O}}_3+2{\mathrm{Si}}_{\mathrm{Si}}^{\mathrm{X}}+{\mathrm{Li}}_2\mathrm{O}\to 2{\mathrm{R}}_{\mathrm{Si}}^{\prime }+2{\mathrm{Li}}_{\mathrm{i}}^{•}+2{\mathrm{SiO}}_2$$

(12)

The results derived from these calculations are summarized in Figure 4. The most promising dopant was the Al³⁺ (see Figure 4). This was due to the ionic radius of Al³⁺ (0.39 Å) matching closely with the ionic radius of Si⁴⁺ (0.26 Å). The solution energy calculated for the Ga³⁺ was higher—only by 0.06 eV—than that calculated for the Al³⁺, suggesting that this dopant is also worth trying experimentally. A gradual increase in the solution energy with increasing ionic radius was observed. The doping of La³⁺ was unfavorable as its solution energy was highly positive (2.19 eV). Previous experimental studies show that Al doping in various Li-containing electrode materials enhances the capacity and electrochemical performance [50,51,52]. In a previous simulation study, it was shown that Al doping on the Si site is the lowest energy process in Li₂MnSiO₄ [53]. It was later confirmed in an experimental study by Duncan et al. [54].

The DFT-optimized structure of Al-doped Li₆SiO₄Cl₂ is shown in Figure 5a. The doped Al formed longer Al–O bonds with nearest neighbor oxygen atoms than Si–O bonds. This was associated with the fact that the ionic radius of Al³⁺ is larger than that of Si⁴⁺. Furthermore, AlO₄ tetrahedral units were slightly distorted upon doping. The DOS plots show that there was a slight increase in the Fermi energy. However, insulating character of this material was unchanged.

Different tetravalent dopants were next substituted on the Si site. The following reaction equation explains the doping process.

$${\mathrm{RO}}_2+{\mathrm{Si}}_{\mathrm{Si}}^{\mathrm{X}}\to {\mathrm{R}}_{\mathrm{Si}}^{\mathrm{X}}+{\mathrm{SiO}}_2$$

(13)

Solution energies calculated for the tetravalent dopants are provided in Figure 6. The energy of solution calculated for the dopant Ge⁴⁺ was −0.01 eV (see Figure 6). This dopant was predicted to be thermodynamically favorable. The suitability of this dopant was due to its ionic radius being closer to that of Si⁴⁺. The second promising dopant was Ti⁴⁺ with a solution energy of 0.59 eV. The other three dopants (Sn⁴⁺, Zr⁴⁺ and Pb⁴⁺) exhibited higher solution energies (>1.70 eV) owing to their ionic radii being larger than the ionic radius of Si⁴⁺. The doping of Pb⁴⁺ was highly unfavorable with a solution energy of 2.15 eV. In previous theoretical studies, the Ge⁴⁺ ion was identified as a candidate dopant ion on the Si site in silicate materials [37,55]

The longer Ge–O bond distances were the result of the larger ionic radius of Ge⁴⁺ (see Figure 7a) although the tetrahedral units were not much distorted. Both Femi energy level and the band gap value were not altered significantly upon Ge doping (see Figure 7b,c).

Finally, the Cl site was considered for doping other halogen atoms according to the following reaction equation.

$$\frac{1}{2}{\mathrm{R}}_2+{\mathrm{Cl}}_{\mathrm{Cl}}^{\mathrm{X}}\to {\mathrm{R}}_{\mathrm{Cl}}^{\mathrm{X}}+\frac{1}{2}{\mathrm{Cl}}_2$$

(14)

Calculated solution energies vs. ionic radii of halogen ions are shown in Figure 8. The smaller ionic radius of F⁻ compared with Cl⁻ reflected in the exothermic solution energy (−1.99 eV) (see Figure 8). Solution energies were positive for both Br⁻ (0.44 eV) and I⁻ (1.37 eV) as their ionic radii are larger than that of Cl⁻. High positive solution energy of I⁻ indicated that heat energy was required for this dopant to overcome the activation energy. Substitution of F⁻ on the Cl or Br site in the antiperovskite Li₂OHR (R = Cl or Br) solid electrolyte material was shown to exhibit excellent electrochemical stability with high voltage [56].

The optimized F-doped structure is shown in Figure 9a. The Li–F bond distances were shorter than the Li–Cl bond distances, indicating the strong ionic bonding between Li⁺ and F⁻ ions. The DOS plots show that the doped configuration was an insulator and the band gap was unaltered upon doping (see Figure 9b,c).

3.4. Formation of Li₆SiO₄Cl₂

Different chemical reaction routes were established to calculate the formation energy of one mole of Li₆SiO₄Cl₂.

Table 4. Calculated reaction energies for the formation of Li₆SiO₄Cl₂.

Reaction Number	Reaction	Reaction Energy (eV)
1	Li4SiO4 + 2 LiCl → Li6SiO4Cl2	−0.12
2	Li2O + Li2SiO3 + 2 LiCl → Li6SiO4Cl2	−0.68
3	2 Li2O + SiO2 + 2 LiCl → Li6SiO4Cl2	−1.98
4	Li2O + Li3OCl + SiO2 + LiCl → Li6SiO4Cl2	−2.09
5	2 Li3OCl + SiO2 → Li6SiO4Cl2	−2.20
6	6 Li + Si + 2 O2 + Cl2 → Li6SiO4Cl2	−30.54

reports the type of reactions and their reaction energies. In all cases, atom economy [57] ($\mathrm{atom}\mathrm{economy}=\frac{\mathrm{mass}\mathrm{of}\mathrm{desired}\mathrm{product}}{\mathrm{total}\mathrm{mass}\mathrm{of}\mathrm{reactants}}\times 100$) was 100%. In the first reaction, Li₄SiO₄ and LiCl were considered as reactants and the reaction was exothermic with a reaction energy of −0.12 eV. The second reaction was even more exothermic, and in this reaction Li₂O, Li₂SiO₃ and LiCl were used as reactants. In the reaction 3, one mole of SiO₂ and one mole of Li₂O were used as reactants instead of one mole of Li₂SiO_3, and the reaction energy was more negative than that calculated in reaction 2. The fourth reaction was slightly more favourable than the third reaction. In this reaction process, one mole of Li₃OCl was used instead of one mole of LiCl and one mole of Li₂O. An additional mole of Li₃OCl was introduced and both Li₂O and LiCl were removed in the fifth reaction. This reaction was more feasible, showing that the Li₃OCl should be a promising ingredient in the preparation of Li₆SiO₄Cl₂. This could partly be due to the exothermic reaction of Li₂O + LiCl → Li₃OCl (−0.22 eV, calculated in this study). Finally, constituent elements were used as reactants in reaction six, to prepare Li₆SiO₄Cl₂. The reaction was hugely negative, meaning that this route is theoretically feasible. Furthermore, this reaction pathway was over an order of magnitude more exothermic than the others listed in the table.

3.5. Li-Incorporation

Thermodynamical stability of Li₆SiO₄Cl₂ upon extra Li (up to three) was examined by relaxing the ion positions and lattice constants of Li-incorporated structures. The possible configurations considered in this study are shown in the ESI (Figures S1–S3). Incorporation energies, Bader charges on the incorporate Li atoms, volume changes and electronic structures of resultant configurations were calculated. The optimized configurations are shown in Figure 10. The calculated incorporation energies were positive, inferring the instability of the incorporated Li atoms (see

Table 5. Incorporation energies, Bader charges on the Li atoms incorporated and change in volume upon successive incorporation.

Reaction	Incorporation Energy (eV)/Li		Bader Charge (|e|)	$\frac{\mathbf{\Delta }\mathit{V}}{\mathit{V}}(\mathbf{\%})$
	Ref: Li Atom	Ref: Li Bulk
Li + Li6SiO4Cl2 → Li•Li6SiO4Cl2	0.47	2.08	+1.00	0.92
2 Li + Li6SiO4Cl2 → 2 Li•Li6SiO4Cl2	0.30	1.91	+1.00 (2)	1.34
3 Li + Li6SiO4Cl2 → 3 Li•Li6SiO4Cl2	0.28	1.89	+1.00 (3)	1.94

). The energies to incorporate Li atoms calculated using bulk Li as chemical potential were highly positive compared with those calculated using Li atom as chemical potential. This was due to the additional energy needed to dissociate bulk Li to form Li atoms. Bader charge analysis showed the loss of electrons from Li atoms forming single positively charged Li⁺ ions. A gradual increase in the volume was observed for each Li incorporation. Such increment in the volume facilitated the incorporation. The Fermi energy level shifted significantly for each incorporation and the resultant configurations became metallic with the loss of electrons from the incorporated Li atoms in the lattice.

4. Conclusions

Ab initio simulations based on the density functional theory were performed to examine the defect properties, doping process, different synthetic routes, and Li-incorporation. We found that the most thermodynamically promising defect process was the O–Cl anti-site defect cluster, indicating that there was a small amount of anion mixing present in this material. The LiCl Schottky was the second most favorable defect. This defect can lead to the loss of LiCl. The Li Frenkel was also a competent defect and its defect energy was 0.06 eV lower than the LiCl Schottky. Na, Ge and F were identified as the most promising isovalent dopants on the Li, Si and Cl sites, respectively. Exerting additional Li-ions was possible by the doping of Al on the Si site. The formation reaction of Li₆SiO₄Cl₂ was examined and the reaction consisting of constituent elements as reactants was highly exothermic. Incorporation of extra Li in this material was not feasible with respect to both Li atom and Li bulk. Furthermore, incorporated Li atoms became Li⁺ ions losing their outer electrons, and there was a gradual expansion in the volume. Exoergic solution energies calculated for Na⁺, Ge⁴⁺ and F⁻ necessitate experimental verification.