Advances in high-capacity Li₂MSiO₄ (M = Mn, Fe, Co, Ni, …) cathode materials for lithium-ion batteries

Abstract

The orthosilicate, Li₂MSiO₄ (M = Mn, Fe, Co, Ni, …), is a competitive cathode material for next generation high-capacity (∼330 mA h g⁻¹ with a possible 2 Li⁺ extraction per formula unit) rechargeable lithium-ion batteries. It is an alternative to other polyanionic compounds because the raw materials of Fe/Mn-containing oxidates and silica are abundant, cost effective, very safe, environmentally friendly, and its Si–O bond is at least as stable as other polyoxyanion groups with an excellent cycling performance. This review highlights the research progress related to the Li₂MSiO₄ cathode materials in terms of the phase transition of the structure, synthetic methods and techniques for improving the electrochemical performance.

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H.-N. Girish

Dr H.-N. Girish received his Bachelor, Master and PhD degree at Uni. of Mysore, Karnataka, India under the supervision of Prof. B. Basavalingu. He is currently working as a postdoctoral researcher in Prof. G.-Q. Shao’s group in the State Key Lab of Adv. Technol. for Mater. Syn. Proc., Wuhan Uni. of Technol., Wuhan, Hubei, China. Present research involves the synthesis, phase transition and electrochemical properties of orthosilicate Li2MSiO4 (M = Mn, Fe, Co, Ni, etc.) cathode materials for rechargeable Li-ion batteries.

G.-Q. Shao

Dr Gang-Qin Shao (G.-Q. Shao), has been a professor at the State Key Lab of Adv. Technol. for Mater. Syn. Proc., Wuhan Uni. of Technol. (WUT), a postdoctoral fellow at Shanghai Inst. Ceram., a PhD and M.Eng. student at WUT, and a B.Eng. student at South China Uni. of Technol., and has been working as a visiting scholar in DMSE, Massachusetts Institute of Technology and in Geballe Lab for Adv. Mater., Stanford Uni., USA, and worked in Yangtze River Sci. Res. Inst., China. Current research interests include: (1) advanced inorganic materials – structure constitution, resolution and refinement; (2) novel functional materials such as rechargeable Li-ion batteries & ferroelectric/magnetoelectric materials.

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A. Introduction

Lithium-ion batteries (LIBs) can effectively store energy in the form of chemicals. Cathode materials have become a hot topic for LIBs. Early cathode materials such as LiCoO₂, LiNiO₂, LiMn₂O₄, LiMnO₂, etc. have limitations related to cost, toxicity to the environment, and inherent chemical and thermal stability in their applications. Extensive research can be found, which has focused on developing suitable promising cathodes^1–6 in the field of polyanionic materials such as phosphates (LiMPO₄, Li₃M₂(PO₄)₃, LiMP₂O₇, M = Mn, Fe, Co, Ni, …, the same below),^7–17 silicates (Li₂MSiO₄, Li₂VOSiO₄),^18–24 fluorophosphates (Li₂MPO₄F, Li₂VPO₄F),²⁵ fluorosilicates (R₂SiF₆, R = K, Na, NH₄⁺, …), sulfates (LiCoOSO₄, LiNiOSO₄, Li₂NiSO₄),^26,27 tavorite fluorosulphates (LiMSO₄F)²⁸ and borates (LiMBO₃).^29–33 Recently the lithium transition-metal (TM) orthosilicate, Li₂MSiO₄ (M = Mn, Fe, Co, Ni, …), has been attracting considerable attention for use as a new generation cathode. It is an alternative to phosphates, or possibly a better option because of its strong Si–O covalent bond (at least as stable as other polyoxyanion groups) resulting in high chemical stability towards electrolytes. It has a high theoretical capacity (∼330 mA h g⁻¹) which can potentially extract more than one lithium ion per formula unit (i.e. 1 Li⁺/f.u.).^{18,23,34–39}

Calculations of the electrochemical potential of lithium insertion or extraction show two processes:

Li₂M²⁺SiO₄ → LiM³⁺SiO₄ + Li⁺ (1)

LiM³⁺SiO₄ → M⁴⁺SiO₄ + Li⁺ (2)

Depending on the TM type, as shown in Fig. 1, the potential in reaction (1) (green square) varies from 3.1 V (for Fe) to 4.1–4.6 V (for Mn, Co and Ni), while that in reaction (2) (red circle) is in the range 4.5–5.1 V. The extraction of the 2^nd Li⁺ requires the application of a high potential above 4.5 V which is a challenge for electrolytes.

Fig. 1 Variation of Li insertion voltage by M²⁺/M³⁺ and M³⁺/M⁴⁺ in Li₂MSiO₄.¹⁸

Li₂MnSiO₄ has more advantages than Li₂FeSiO₄ with regards to cell safety, ease of preparation and cost effectiveness. The Mn³⁺ ↔ Mn⁴⁺ conversion provides a higher potential than Mn²⁺/Mn³⁺ and Fe²⁺/Fe³⁺ (it is yet to be clarified whether this involved an Fe³⁺/Fe⁴⁺ redox couple which might just be an electrolyte degradation^18,38,40,41). During oxidation, the short Mn–O bonds in the MnO₄ tetrahedron decrease in correlation with the increase in the Mn valance state, while the long Mn–O bonds remain unchanged.⁴² In practical applications, Li₂MnSiO₄ is limited as a cathode by its low electronic conductivity of ∼5 × 10⁻¹⁶ S cm⁻¹ at room temperature (RT) (and ∼3 × 10⁻¹⁴ S cm⁻¹ at 60 °C), which is 5–6 orders of magnitude smaller than that of LiFePO₄ at RT. The preparation of pure Li₂MnSiO₄ even at temperatures up to 800 °C is not trivial due to the possible presence of mixed phases and impurities such as MnO, MnSiO₃, Li₂SiO₃, etc.^43–46

Co- or Ni-containing salts for preparing Li₂CoSiO₄ and Li₂NiSiO₄ are commercially expensive and environmentally toxic. Li₂NiSiO₄, with the lowest band-gap, does not have the shortcomings of Mn derivatives¹⁸ but the delithiation potentials (>4.6 V) of Li₂CoSiO₄ and Li₂NiSiO₄ are too high to make the present electrolytes suitable Li⁺ diffusion media. In the same way, their kinetic properties make preparation difficult and cause poor delithiation–lithiation during the charge–discharge process.^{18,36,47–49}

Otherwise, defects such as poor reversibility that are associated with Li₂MSiO₄ have become a huge obstacle to their extensive application in electric vehicles (EV) and plug-in hybrid electric vehicles (PHEV).^50–53 Various methods have been developed to improve their electrochemical properties, such as controlling the particle size and morphology (e.g., nanosheet/rod-like particles),^45,54–57 cation substitution (Mg²⁺, Al³⁺, Cr²⁺, etc.)^56,58–65 and surface-coating with conductive materials.^66–80 For surface-coating methods, an appropriate carbon coating has been widely used,^81–84 because it not only improves the electronic conductivity but also reduces the particle size, which are favourable for enhancing the electrochemical properties.

This review will explain the polymorphic structure of Li₂MSiO₄ and its phase transitions and summarize progress in its research, methods for its synthesis and techniques for improving its electrochemical performance.

B. Polymorphic structure of Li₂MSiO₄ and its phase transitions

The orthosilicate (Li₂MSiO₄) represents a family having an iso-structure of tetrahedral lithiophosphate (Li₃PO₄). The crystal structure of Li₂MSiO₄ consists of distorted hexagonal close-packing of oxygen ions with half tetrahedral sites occupied by Li, M (M = Mn, Fe, Co, Ni, …) and Si cations, so that face sharing between pairs of tetrahedral sites is avoided. These tetrahedral structures exhibit a rich polymorphism.

Li₃PO₄ crystallizes in the “α-phase formed in low temperature”, “β-phase in high temperature” and “γ-phase above 1170 °C”.^45,85 Several temperature/pressure-dependent polymorphs of Li₂MSiO₄, namely high temperature monoclinic γ_o/γ_s (P2₁/n) and orthorhombic γ_II (Pmna, Pmnb, Cmma, γ-Li₃PO₄), and low temperature orthorhombic β_I (Pbn2₁, Pna2₁, β-Li₃PO₄)/β_II (Pmn2₁, β-Li₃PO₄), can be stabilised/influenced significantly by the synthetic conditions and also the electrochemical cyclic behavior.⁸⁶ These phases may be quenched to RT and exhibit long-term stability. Metastable monoclinic (P_n) Li₂MnSiO₄ prepared by ion-exchange from Na₂MnSiO₄ was also reported.⁸³ Different phase structures of Li₂MSiO₄ polymorphs are shown in Fig. 2 and discussed below.^{38,45,83,86–92}

Fig. 2 Structures of Li₂MSiO₄: (a) orthorhombic β_II (Pmn2₁); (b) orthorhombic β₁ (Pbn2₁, Pna2₁); (c) orthorhombic γ_II (Pmna, Pmnb, Cmma); (d) monoclinic γ_s (P2₁/n, P2₁); (e) monoclinic γ₀ (P2₁/n); (f) monoclinic P_n.^{38,45,83,86–92} SiO₄ (yellow); MO₄ (brown/green); LiO₄ (blue/orange); light and dark blue tetrahedra represent crystallographically distinct Li sites.

(a) Orthorhombic β_II structure (Pmn2₁) and modified-/inverse-β_II structure (Pmn2₁-I): all the tetrahedra point in the same direction, perpendicular to the close-packed planes, and share only corners with each other. Chains of LiO₄ are along the a-axis and parallel to chains of alternating MO₄ and SiO₄ (Fig. 2a).⁴⁵

In the modified-/inverse-β_II structure (e.g. the cycled polymorph of Li₂FeSiO₄), all tetrahedra point in the same direction along the c-axis/b-axis and are linked only by corner-sharing. The SiO₄ tetrahedra are isolated from each other, sharing corners with LiO₄ and (Li/Fe)O₄ tetrahedra.^86,90,91 This the more stable configuration at low temperature.

(b) Orthorhombic β_I structure (Pbn2₁, Pna2₁): all tetrahedra point in same direction with chains of alternating LiO₄ and MO₄ tetrahedra along the a-axis, parallel to chains of alternating LiO₄ and SiO₄ tetrahedra (Fig. 2b).^{38,87–90,92}

Quoirin et al. prepared an orthorhombic Li₂FeSiO₄ (Pna2₁, β_I) by precipitation at 100 °C in air which was optionally in the space group of Pccn. It might be an intermediate between the orthorhombic Pmn2₁ (β_II) and monoclinic P2₁/n (γ_s) phases.^87–89

(c) Orthorhombic γ_II structure (Pmna, Pmnb, Cmma): the tetrahedra are arranged in groups of three with the central tetrahedron pointing in the opposite direction to the outer two, with which it shares edges; the group of 3 edge-sharing tetrahedra consist of the sequence Li–M–Li (Fig. 2c and the inset).⁴⁵

(d) Monoclinic γ_s structure (P2₁/n, previously designated P2₁ (ref. 93–95)): half of the tetrahedra point in reverse directions, containing pairs of LiO₄/MO₄ and LiO₄/LiO₄ edge-sharing tetrahedra (Fig. 2d and the inset).⁴⁵

(e) Monoclinic γ₀ structure (P2₁/n): similar to the γ_II structure, the group of 3 edge-sharing tetrahedra consist of the sequence Li–Li–M (Fig. 2e).⁴⁵

(f) Monoclinic structure (P_n): similar to the orthorhombic β_II structure (Pmn2₁), it is a metastable form of Li₂MnSiO₄ and can convert to β_II above 370 °C. Here the MnO₄ and SiO₄ tetrahedra point in the same direction parallel to the c-axis (Fig. 2f).⁸³

Fig. 3 shows a schematic diagram of the temperature/pressure-dependent polymorphs of Li₂MSiO₄ with phase transitions. Table 1 shows the reported lattice parameters of Li₂MSiO₄ polymorphs (M = Fe, Mn, Co, Ni).

Fig. 3 Schematic diagram of temperature/pressure-dependent polymorphs of Li₂MSiO₄ with phase transitions.

Table 1 The reported lattice parameters of Li₂MSiO₄ polymorphs (M = Fe, Mn, Co, Ni)

Li2MSiO4	x	Space group	Crystal system	a/Å	b/Å	c/Å	α/°	β/°	γ/°	Refs.
Li2FeSiO4	—	Pbn21 (βI)	Orthorhombic	6.2652(1)	10.8121(1)	4.9504(1)	90	90	90	88 and 89
	—	Pna21 (βI)	Orthorhombic	10.724(5)	6.260(7)	4.958(6)	90	90	90	87–89
	—	Pmn21 (βII)	Orthorhombic	6.2709(8)	5.3382(7)	4.9651(7)	90	90	90	88 and 89
	—	Pmn21 (βII)	Orthorhombic	6.2661	5.3295	5.0148	90	90	90	23
	—	Pmn21 (βII)	Orthorhombic	6.271(1)	5.336(1)	4.9607(9)	90	90	90	37
	—	Pmn21 (βII)	Orthorhombic	6.3246	5.3817	4.9967	90	90	90	18
	—	Pmn21 (βII)	Orthorhombic	6.313	5.393	4.979	90	90	90	39
	—	Pmnb (γII)	Orthorhombic	6.2853(5)	10.6592(8)	5.0367(4)	90	90	90	103
	—	Pnma (γII)	Orthorhombic	10.656	6.285	5.038	90	90	90	88
	—	Cmma (γII)	Orthorhombic	10.664	12.540	5.022	90	90	90	88
	—	P21/n (γs)	Monoclinic	8.2253(5)	5.0220(1)	8.2381(4)	90	99.230	90	93
	—	P21/n (γs)	Monoclinic	6.2835(7)	10.6572(1)	5.0386(5)	90	89.941	90	103
	—	P21 (γs)	Monoclinic	8.2278(1)	5.0204(6)	8.2295(1)	90	99.231	90	95
Li2MnSiO4	—	Pmn21 (βII)	Orthorhombic	6.3109(9)	5.3800(9)	4.9662(8)	90	90	90	37
	—	Pmn21 (βII)	Orthorhombic	6.308(3)	5.377(3)	4.988(9)	90	90	90	107
	—	Pmn21 (βII)	Orthorhombic	6.3666	5.4329	4.0368	90	90	90	18
	—	Pmn21 (βII)	Orthorhombic	6.308(9)	5.385(9)	4.999(6)	90	90	90	108
	—	Pmn21 (βII)	Orthorhombic	6.3064	5.3893	4.9715	90	90	90	109
	—	Pmn21 (βII)	Orthorhombic	6.3141	5.3702	4.9650	90	90	90	44
	—	Pmnb (γII)	Orthorhombic	6.3126	10.7657	5.0118	90	90	90	44
	—	P21/n (γs)	Monoclinic	6.3368(1)	10.9146(2)	5.0730(1)	90	90.98	90	106
	—	P21/n (γs)	Monoclinic	6.3363(5)	10.8950(7)	5.07506(3)	90	90.99	90	44
	—	Pn (metastable)	Monoclinic	6.593(5)	5.402(1)	5.090(2)	90	89.7(3)	90	83
Li2CoSiO4	—	Pbn21 (βI)	Orthorhombic	6.253	10.685	4.929	90	90	90	92
	—	Pmn21 (βII)	Orthorhombic	6.2599	10.6892	4.9287	90	90	90	85
	—	Pmn21 (βII)	Orthorhombic	6.159	5.440	4.988	90	90	90	110
	—	Pmn21 (βII)	Orthorhombic	6.287(6)	5.353(1)	4.937(1)	90	90	90	47
	—	Pmn21 (βII)	Orthorhombic	6.267(9)	5.370(8)	4.939(4)	90	90	90	47
	—	Pmn21 (βII)	Orthorhombic	6.2015	5.4378	4.9909	90	90	90	18
	—	Pmnb (γII)	Orthorhombic	6.20	10.72	5.03	90	90	90	85
	—	P21/n (γo)	Monoclinic	6.284	10.686	5.018	90	90.60	90	43
Li2NiSiO4	—	Pmn21 (βII)	Orthorhombic	6.2938	5.3702	4.9137	90	90	90	18
	—	Pmn21 (βII)	Orthorhombic	6.274	5.365	4.921	90	90	90	111
	—	Pmn21 (βII)	Orthorhombic	6.290	5.299	4.856	90	90	90	112
Li2(Fe1−xMnx)SiO4 (x = 0, 0.3, 0.5, 0.7, 1)	x = 0	Pmn21 (βII)	Orthorhombic	6.2512	5.3425	5.0188	90	90	90	113
	x = 0.3			6.2117	5.381	4.975
	x = 0.5			6.2581	5.3747	5.006
	x = 0.7			6.2851	5.381	5.0012
	x = 1			6.3099	5.394	5.0192
Li2(Mn1−xFex)SiO4/C (x = 0, 0.2, 0.5, 0.8)	x = 0	Pmn21 (βII)	Orthorhombic	—	—	—	90	90	90	109
	x = 0.2			—	—	—
	x = 0.5			—	—	—
	x = 0.8			—	—	—
Li2(Fe1−xMnx)SiO4 (x = 0, 0.2, 0.5, 1)	x = 0	P21/n (γs)	Monoclinic	8.244(8)	5.004(4)	8.244(8)	90	99.25	90	102
	x = 0.2	P21/n (γs)	Monoclinic	8.264(7)	5.020(7)	8.264(4)	90	99.00	90
		Pmnb (γII)	Orthorhombic	6.591(6)	9.462(7)	4.988(1)	90	90	90
	x = 0.5	Pmn21 (βII)	Orthorhombic	6.284(0)	5.391(9)	5.015(1)	90	90	90
		P21/n (γs)	Monoclinic	6.986(8)	8.030(9)	5.990(4)	90	95.19	90
	x = 1	Pmn21 (βII)	Orthorhombic	6.299(9)	5.381(5)	4.974(2)	90	90	90
		Pmnb (γII)	Orthorhombic	5.929(5)	10.622(0)	4.961(1)	90	90	90
		P21/n (γs)	Monoclinic	6.305(7)	11.098(7)	5.114(9)	90	92.44	90
Li2(Fe1−xMnx)SiO4/C (x = 0, 0.05, 0.1, 0.2, 0.3)	—	Pmn21 (βII)	Orthorhombic	—	—	—	90	90	90	114
Li2(Fe0.5Mn0.5)SiO4	—	Pmnb (γII)	Orthorhombic	—	—	—	90	90	90	115
Li2(Fe0.5Mn0.5)SiO4/C	—	Pmn21 (βII)	Orthorhombic	—	—	—	90	90	90	106

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The differences among these polymorph structures mainly depend on the different arrangements of the cationic tetrahedra and the local order (order probes such as solid-state Magic Angle Spinning Nuclear Magnetic Resonance (MAS NMR) are very useful for identifying the obtained polymorphs and also their mixtures).^96,97

B.1. Synthetic conditions for phase formation and phase transition in Li₂MSiO₄

The structures of Li₂MSiO₄ materials obtained from different synthetic methods further affect the lithium intercalation behavior.¹¹⁶ Among the various polymorphs, Li₂MSiO₄ with P2₁ symmetry has higher electronic conductivity⁴⁸ but both high/low temperature orthorhombic structures are more stable than the monoclinic ones.⁴⁴

Orthorhombic β_II-Li₂MSiO₄ (Pmn2₁) can be formed at low temperature (150–200 °C) via hydrothermal methods,^{100,117–120} and at 600–800 °C via solid-state,^{37,73,81,99,121–128} sol–gel,^{41,57,78,84,108,129–138} spray pyrolysis,^139,140 microwave¹⁴¹ and combustion¹⁴² methods, etc. It is very stable and easy to synthesize via all these techniques. For example, β_II-Li₂CoSiO₄ was obtained via a hydrothermal technique at 150 °C and treatment at 700 °C,¹⁴³ via a combustion route at 950 °C (ref. 144) and via a supercritical technique at 350 °C/38 MPa.¹⁴⁵

Orthorhombic β_I-Li₂MSiO₄ (Pbn2₁, Yamaguchi et al. named it β_Il-Pbn2₁ previously⁹²) is formed at a higher temperature than β_Il-Pmn2₁ and little data have been reported for this phase.^143–145 To the best of our knowledge, this phase is not yet reported for Li₂MnSiO₄.

Monoclinic γ_s-Li₂MSiO₄ (P2₁/n) is formed at a higher temperature than the β_I- and β_Il-phases in the temperature range 650–900 °C via solid-state,^{43,65,79,82,105} sol–gel^{56,61,63,129,146} and combustion^43,147,148 methods. When the temperature was further increased to 800–1000 °C, an orthorhombic γ_II-phase (Cmma,^87,88 Pnma¹⁴⁹ or Pmnb^{44,85,87,88,100,103,129,150}) resulted.

The β-phase is more stable than the γ-phase because the latter has a larger volume and requires a higher temperature for preparation and quenching.⁴⁴

Li₂FeSiO₄ (LFS). Li₂FeSiO₄ crystallises in the orthorhombic β_II (Pmn2₁), orthorhombic γ_II (Pmnb) and monoclinic γ_s (P2₁/n) phases. The inverse-β_II structure (Pmn2₁-I, β_II) is commonly formed at low temperatures. The monoclinic and orthorhombic structures stabilized depending on the synthetic method, different temperatures and carbon-coating technique.^42,151

Nyten et al. obtained pure orthorhombic β_II-Li₂FeSiO₄ (Pmn2₁) via a solid-state method.^{23,99,152,153} Quoirin et al. synthesized orthorhombic β_I-Li₂FeSiO₄ (Pna2₁) via a wet chemical method at 100 °C, and orthorhombic γ_II-Li₂FeSiO₄ at 800 °C (Cmma)/900 °C (Pnma) via a solid-state method.^87–89 Nishimura et al. prepared monoclinic γ_s-Li₂FeSiO₄ at 800 °C (P2₁/n).^93–95 Sirisopanaporn et al. obtained orthorhombic γ_II-Li₂FeSiO₄ (Pmnb) at 900 °C that differs from the monoclinic γ_s-phase obtained by quenching from 800 °C.^93,103 Bini et al. synthesized pure monoclinic γ_s-Li₂FeSiO₄ (P2₁/n) at 650 °C and pure orthorhombic γ_II-Li₂FeSiO₄ (Pmnb) at 900 °C via a sol–gel method.¹²⁹

At 580 °C, complete phase transformation (starting at 400 °C) from orthorhombic (Pmn2₁, β_II) to monoclinic (P2₁/n, γ_s) was observed and it was maintained up to 820 °C. Above 900 °C, it phase transformed again into an orthorhombic phase (Pmna, γ_II). This involved a statistical distribution of Li and Fe over the three (2/3 Li and 1/3 Fe) cation sites. By quenching the orthorhombic phase (Pmna, γ_II) from 900 °C, the monoclinic phase (P2₁/n, γ_s) was fully regained at 600 °C and maintained down to RT.¹⁰¹ Nanorods of orthorhombic Li₂FeSiO₄ (Pmn2₁, β_II) were prepared at 200 °C over 6 days and phase transformed to a monoclinic phase (P2₁/n, γ_s) by heating at 600 °C.⁵⁷

Li₂MnSiO₄ (LMS). There are four polymorphs of Li₂MnSiO₄ that form at ambient pressure. Two of them are orthorhombic β_II and γ_II (i.e., Pmn2₁ and Pmnb). Another two are monoclinic P2₁ (γ_s) and P_n.

Politaev et al. synthesized monoclinic (P2₁/n) Li₂MnSiO₄ at 950–1050 °C via a solid-state method.¹⁰⁶ As mentioned before, the monoclinic P_n phase would transform to a stable orthorhombic phase (Pmn2₁) above 370 °C.⁸³ By rate-cooling, a reversible phase transition can be observed in Li₂MnSiO₄. The sample (Pmnb) prepared at 800 °C via a solid-state method transformed to a P2₁/n phase when it cooled to RT at 200 K min⁻¹, but the sample prepared at 700 °C formed the Pmn2₁ phase. This strongly indicated that the Pmnb phase transformed to orthorhombic disordered wurtzite-type structures.⁶¹

Bini et al. synthesized pure orthorhombic β_II-Li₂MnSiO₄ (Pmn2₁) at 650 °C and mixed phases of orthorhombic γ_II (Pmnb)/monoclinic γ_s (P2₁/n) at 900 °C via a sol–gel method.¹²⁹

By a hydrothermal technique and heating at 700 °C, orthorhombic Li₂MnSiO₄ (Pmn2₁, β_II) was synthesized. Above 750 °C, it changed into a monoclinic (P2₁/n, γ_s) phase due to discontinuity in the c lattice parameter.¹⁰¹ By high pressure and high temperature processes, phase transitions can be observed with minimal impurities. The mixed phase of Li₂MnSiO₄ polymorphs, i.e. Pmnb and P2₁/n (63 : 13 w/w) was prepared at 900 °C over 10 h by slowly heating (2 °C min⁻¹) with small impurities of MnO and Li₂SiO₃. At 60 kbar and 600 °C, the mixed phase transformed to a Pmn2₁ phase, being stable at 80 kbar and 900 °C with a new impurity of Mn₂SiO₄.⁴³

The low-temperature orthorhombic forms are more stable (due to their large volume) than the monoclinic form, which can only be prepared above 900 °C.^44,97 By increasing the process pressure/temperature, polymorphic transformations (Pmn2₁ → Pmnb → P2₁/n) will take place.^44,144

Li₂CoSiO₄ (LCS). By varying the cooling rate, different polymorphs can form in Li₂CoSiO₄. Rapid cooling (0.5 °C min⁻¹) from above 1000 °C resulted in the formation of Pmnb (γ_II), while slow cooling (0.1 °C min⁻¹) caused the formation of P2₁/n (γ_o). Pmn2₁ (β_II) was formed by static heating of any other phase at 640–800 °C followed by quenching, and Pbn2₁ (β_I) was formed by slow cooling of Pmn2₁ (β_II).⁸⁵

Monoclinic Li₂CoSiO₄ (P2₁/n, γ_o) was prepared at 950 °C and could transform (at 900 °C) into orthorhombic Pmn2₁ (β_Il) and Pbn2₁ (β_I) phases under 40/60 kbar, respectively. Both cell volumes are smaller than the monoclinic one.⁴³ Orthorhombic Li₂CoSiO₄ (Pmn2₁, β_II) was synthesized hydrothermally (150 °C/72 h). It transformed to the orthorhombic (Pbn2₁, β_I) phase when heated at 700 °C over 2 h and to a monoclinic (P2₁/n, γ_o) phase at 1100 °C over 2 h.³⁶ Otherwise, a considerably slow phase transition from β_II ↔ γ_II at 170 °C has been determined by DTA studies.⁸⁵

Li₂NiSiO₄ (LNS). Similar to the others in the family of Li₂MSiO₄, Li₂NiSiO₄ was most frequently observed to be an orthorhombic structure (Pmn2₁, β_II). DFT calculations predict that Li₂NiSiO₄ has very high deintercalation voltages (4.5 V for Ni²⁺/Ni³⁺ and 5.2 V for Ni³⁺/Ni⁴⁺). This causes great difficulty for delithiation–lithiation during the charge–discharge process, even though the Li₂NiSiO₄ powder had been successfully synthesized.^18,111,112

Li₂(Mn_1−xFe_x)SiO₄/Li₂(Fe_1−xMn_x)SiO₄. This polymorph reflects the behavior of the end-members. For Fe-rich compositions, there is a gradual transition from Fe-like behavior to Mn-like behavior with increasing Mn-content.⁹⁸ For Mn-rich compositions, there is a gradual transition from Mn-like behavior to Fe-like behavior with increasing Fe-content.^102,109,113

Li₂(Fe_1−xMn_x)SiO₄ (x = 0, 0.2, 0.5, 1) was prepared via a modified sol–gel route using polyvinylpyrrolidone (PVP) as the chelating agent and carbon source. A pure monoclinic phase (P2₁/n) was obtained when x = 0. Monoclinic (P2₁/n) and orthorhombic (Pmnb) mixed phases were obtained when x = 0.2. Major orthorhombic (Pmn2₁) and minor monoclinic (P2₁/n) mixed phases were obtained when x = 0.5. Major orthorhombic (Pmn2₁), minor orthorhombic (Pmnb) and minor monoclinic (P2₁/n) mixed phases were obtained when x = 1.¹⁰²

However, Li₂(Fe_1−xMn_x)SiO₄ (x = 0, 0.3, 0.5, 0.7, 1) obtained via a citric acid assisted sol–gel technique could be indexed to the orthorhombic phase (Pmn2₁), and the lattice parameters were similar.¹¹³ All samples of Li₂(Fe_1−xMn_x)SiO₄/C (x = 0, 0.05, 0.1, 0.2, 0.3) obtained via a solution route exhibited an orthorhombic phase (Pmn2₁) with a small amount of impurities such as Fe₃O₄ and Li₄SiO₄.¹¹⁴

Li₂(Mn_xFe_1−x)SiO₄ (x = 0, 0.2, 0.5, 0.8) was prepared via a combination of spray pyrolysis at 400 °C and wet ball milling followed by annealing at 600 °C in N₂ for 4 h. They exhibited a major orthorhombic phase (Pmn2₁) and small amounts of impurities such as FeO and SiO₂ were found in materials with high Fe contents of 0.5 and 0.8.¹⁰⁹

The thermal behavior of Li₂(Fe_0.5Mn_0.5)SiO₄ obtained via a sol–gel route, from RT to 950 °C, exhibited a major Pmnb phase (determined by in situ XRD) over the explored temperature range. A pure and stable Pmnb phase was formed at 950 °C.¹¹⁵

Li₂(Fe_0.5Mn_0.5)SiO₄/C was prepared using the precursors of Li₂FeSiO₄ and Li₂MnSiO₄ via a sol–gel route followed by heating at 600 °C for 10 h. It exhibited an orthorhombic phase (Pmn2₁) but the atomic ratio was not maintained.¹⁰⁶

B.2. Phase transition while charging/discharging in Li₂MSiO₄

The rich polymorphism in Li₂MSiO₂ with the associated small transition energies is the main factor of the long-term cyclability. The phase transformation has structural, cyclic rate, applied current density, temperature and pressure dependence.^43,100 The extraction/insertion of the 1^st lithium out/into the host structure corresponds to the reaction Li₂M²⁺SiO₄ → LiM³⁺SiO₄ + Li⁺ + e⁻ and is accompanied by the gradual changes of the local environment of the Li nuclei from the initial Li⁺–O–M²⁺ to Li⁺–O–M³⁺ bonds upon battery discharging. During discharge, phase transitions can be observed with the decrease in the crystal size upon cycling. The residual M³⁺ could explain the differences in the observed NMR spectra between the fully discharged sate and the pristine state.¹⁰² The influence of variations in the arrangements of MO₄, in terms of the orientation, size and distortion, on the equilibrium potential were measured during the first oxidation of M²⁺ into M³⁺ for all polymorphs. The stronger M–O bonds result in a higher splitting energy between the bonding and antibonding states, which reduces the M²⁺/M³⁺ redox potential vs. Li⁺/Li⁰.^100,154 An irreversible phase transformation was observed during the first charge, which accompanied a characteristic potential drop.^23,99,104 XRD studies and theoretical calculations also suggested structural rearrangements involving the interchange of some M and Li sites.⁹⁹ The applied current density is important to the kinetic phase transformation in the galvanostatic measurement.^39,100

When the lithium ions were fully extracted from Li_xMSiO₄ (M = Mn, Fe, Co, Ni, …), a larger volume expansion was found in the Fe-/Co-/Ni-systems compared to the Mn-system.¹¹¹

Based on the structural characterization of the Pmn2₁-delithiated materials, Thomas et al. suggested that the voltage shift was due to a structural transformation of the host compound. Armstrong et al. suggested a Li₂FeSiO₄ phase transition from the monoclinic (P2₁) polymorph to the orthorhombic (Pmn2₁-I, inverse-β_II) phase at 50 °C and at C/16 (1C = 160 mA g⁻¹).¹⁷⁴ Chen et al. confirmed the same phase transition happened at RT and at C/20.¹⁰² But the monoclinic phase could be preserved even after cycling.¹⁵⁵

Nanocrystalline Li₂(Fe_1−xMn_x)SiO₄/C (x = 0, 0.2, 0.5, 1) powders were tested electrochemically, and characterized using XRD, ⁷Li MAS NMR spectroscopy, ⁵⁷Fe Mössbauer spectroscopy and in situ XAS to study the structural evolution. Results showed the materials exhibited partially reversible structural changes upon cycling. Amorphization and structural rearrangements from the initial P2₁/n polymorph to Pmn2₁ occurred during the 1^st charge–discharge cycle. Pmn2₁-Li₂FeSiO₄ exhibited a stable cycling performance over the subsequent 100 cycles due to the Fe²⁺/Fe³⁺ process. A competitive redox reaction between the Mn and Fe species was deduced for Li₂Fe_0.8Mn_0.2SiO₄ and Li₂Fe_0.5Mn_0.5SiO₄. A high conversion rate existed in the 1^st charge from Fe²⁺/Mn²⁺ to Fe³⁺/Mn³⁺, while the presence of a small fraction (<7.5 mol%) of cations with higher oxidation states (Fe⁴⁺/Mn⁴⁺) could not be excluded.¹⁰²

C. Synthetic methods

Table 2 gives a summary of synthetic methods, particle size, morphology, experimental conditions and physicochemical/electrochemical properties for Li₂MSiO₄ (M = Fe, Mn, Co). A high performance is observed in Li₂FeSiO₄ but it is yet to be clarified whether an Fe³⁺/Fe⁴⁺ couple could work.^18,38,40,41 Besides Li₂FeSiO₄, the highest initial discharge of 313 mA h g⁻¹ (0.05 C, nearly 2 Li⁺ extracted, 76.7% @ 20 cycles) was attained using PEDOT-coated Li₂MnSiO₄ at 40 °C, which was prepared via a supercritical technique¹⁵⁶ and the second highest initial discharge of 253.4 mA h g⁻¹ (0.03 C) was attained using Li₂MnSiO₄/C which was synthesized via a sol–gel technique.¹³³

Table 2 Synthetic methods, conditions and physicochemical/electrochemical properties of Li₂MSiO₄ (M = Fe, Mn, Co)

Methods	Temp./time/pressure (°C/h/MPa)	Morphology	Size (nm)	Specific capacity (mA h g^−1)/current density	Retention (%)/cycles	Refs.
Solid-state	650/10/—	Agglomerate (Li2FeSiO4/C)	—	144.8/0.1	94.27/13	188
	800/12/—	Agglomerate (Li2MnSiO4/C)	30–80	129/0.06	—/10	128
	650/8/—	Agglomerate (LiFe0.9Mn0.1SiO4/C)	—	158.1/0.03	94.3/30	114
	900/6/—	Agglomerate (Li2FeSiO4/C/rGO)	28.5	191.6/0.1	91.3/60	127
	700/10/—	Agglomerate (Li1.95FeSiO4/C/CNTs)	30–70	148/0.2	99.2/100	159
	700/8/—	Spherical (Li2MnSiO4/C/graphene)	50	215.3/0.05	81.3/40	125
	650/10/—	Agglomerate (LiFe0.97Co0.03SiO4/C)	100–500	199/3.0	71.6/100	65
	700/15/—	Rounded (Li2FeSiO4/C)	200	102/—	91.8/10	126
	700/10/—	Agglomerate (Li1.95FeSiO4/C)	100	142/1.0/100	95.1/100	122
	650/10/—	Agglomerate (LiFe0.95V0.05SiO4/C)	7–10	220.4/0.1	78.7/50	161
	700/10/—	Agglomerate (Li2MnSiO4/C)	13.4–23.3	201.8/—	73.5/15	163
Sol–gel	800/—/—	Spherical (Li2CoSiO4)	460	32/0.02	—/2	164
	700/5/—	Spherical (Li2MnSiO4)	25–30	161/0.05	77.6/50	137
	700/10/—	Irregular (Li2MnSiO4/C)	20–30	240/0.02	45.4/30	130
	700/—/—	Spherical (Li2MnSiO4)	25–30	113/—	—/15	138
	650/10/—	Mesopores (Li2FeSiO4/C)	20–30	163/0.06	96/200	166
	680/10/—	Mesopores (Li2MnSiO4/C)	10–20	164.2/0.06	80/60	136
	650/10/—	Agglomerate (Li1.9Na0.1MnSiO4/C)	30	175/0.01	45/40	167
	650/10/—	Agglomerate LiFeSi0.9V0.1O4/C)	100–200	159/0.06	90/30	63
	700/12/—	Agglomerate (LiFe0.97Mg0.03SiO4)	100	153.2/0.06	98.6/50	61
	700/10/—	Agglomerate (LiMn0.8Fe0.2SiO4/C)	15–30	224/0.05	63.8/50	189
	700/10/—	Honeycomb (Li1.8MnSiO4)	6–8	147.1/0.03	83.3/25	168
	600/10/—	Aggregated (Li2FeSiO4/C-NHT)	20–30	195.5/0.1	—/50	190
	700/10/—	Agglomerate (Li0.5Mn0.5SiO4/C)	220	230.1/0.1	70.4/20	191
	700/12/—	Core–shell (Li2.05Mn0.95P0.05Si0.95O4)	30–60	170/0.06	—/60	170
	700/12/—	Agglomerate (LiFe0.97Zn0.03SiO4)	100	128/0.1	97.5/55	169
	650/10/—	Mesopores (Li2FeSiO4@CMK-3)	50	160/0.1	—/80	192
	600/10/—	Agglomerate (Li2FeSiO4/C)	15	187/0.1	—/50	193
	700/10/—	Agglomerate (Li2.05Fe0.95P0.05Si0.95O4/C)	50–100	213/0.3	—/50	194
	650/6/—	Spherical (Li2MnSiO4/C)	20–50	253.4/0.03	77/20	133
	600/10/—	Nanowires (Li2MnSiO4/C/V2O5)	30	277.0/0.1	—/50	195
	700/12/—	Nanofibers (LMn0.94Cr0.06SiO4/C)	70–150	295/—	65.8/50	64
	650/10/—	Agglomerate (LiMn0.09Ti0.01SiO4)	∼50	211/0.01	—/50	196
Hydrothermal/solvothermal/supercritical	150/92/—	Flower (Li2FeSiO4)	20–50	130/0.1	—/30	117
	700/—/—	Flower (Li2MnSiO4/C)	—	100/0.05	—/40	162
	550/5/—	Spherical (Li2MnSiO4)	644	177/—	32/20	173
	650/10/—	Hollow spheres (Li2FeSiO4)	0.5–2 μm	152/0.05	—/100	178
	600/6/—	Hot-dog (Li2FeSiO4)	200	150/—	—/50	180
	900/4/—	Hollow spheres (Li2CoSiO4/C)	300–400	33/—	—/50	49
	600/5/—	Agglomerate (Al–LiFe0.5Mn0.5SiO4/C)	20–100	216/0.01	—/20	62
	200/72/—	Spherical (Li2FeSiO4/C)	∼20	136/0.2	96.1/100	118
	700/10/—	Flower (Li2MnSiO4)	20	226/0.5	—/10	116
	650/10/—	Shuttle-like (Li2MnSiO4/C)	0.4–0.5 μm	206/0.05	—/50	132
	180/192/—	Hierarchical shuttle (Li2FeSiO4)	31.8	159.4/0.01	97.5/20	120
	600/2/—	Nanorods (Li2FeSiO4/C)	80–100	155/0.06	—/50	175
	650/10/—	Spherical (Li2FeSiO4/C)	18.3	211.3/0.1	97.7/1000	176 and 181
	700/10/—	Plates (LiMn0.8Ni0.2SiO4)	100	100/0.2	—/10	179
	600/10/—	3D porous hierarchical (Li2FeSiO4/C)	∼60	170/0.1	—/20	57
	600/06/—	Nanorods (Li2FeSiO4/graphene)	10–25	306.4/0.3	95/240	97
	700/10/—	Agglomerate (Li2MnSiO4@Ni)/C	20–50	274.5/0.05	—/20	185
	700/5/—	Spherical (Li2MnSiO4)	15–500	260/0.1	—/5	177
	380/30 min/38	Rod (Li2FeSiO4)	20–80	177/0.01	—/25	145
	350/30 min/38	Spherical (Li2CoSiO4)	20–15	107/0.1	—/4	156
	300/5 min/38	Spherical (Li2MnSiO4)	5–20	313/0.05	76.7/20	197
	400/4 min/38	Hierarchical nano (Li2MnSiO4)	4–5	291/0.05	—/50	198
Microwave	650/6/—	Nanosphere (Li2FeSiO4/C)	∼20	148/0.05	—/20	171
	650/6/—	Agglomerate (Li2MnSiO4/C)	∼20	210/0.05	—/20	171
	700/20/—	Agglomerate (Li2FeSiO4)	—	116.2/0.05	—/10	183
	700/20/—	Agglomerate (Li2FeSiO4)	—	116.9/0.05	—/10	184
Spray pyrolysis/combustion/hydrochemical	800/4/—	Agglomerate (Li2MnSiO4/C)	∼50	184/0.05	—/20	139
	700/10/—	Spherical (Li2FeSiO4/C)	1–5 μm	123/1.0	98.4/300	199
	600/4/—	Spherical (Li2FeSiO4/C)	65	154/0.05	—/70	200
	600/4/—	Agglomerate (LiFe0.5Mn0.5SiO4/C)	65	149/1.0	—/50	109
	700/12/—	Agglomerate (Li2FeSiO4/C/C-nanosphere)	∼200	164.7/0.1	98.4/60	187
	650/10/—	Porous (Li2FeSiO4/C)	28	135/0.06	—/—	148
	800/10/—	Agglomerate (Li2FeSiO4/C)	29	130/0.06	—/50	147
	700/10/—	Agglomerate (Li2MnSiO4/C)	50	164/0.01	—/20	142

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C.1. Solid-state method

A high temperature is applied during this process. It appears to be quite simple, but cannot be used for all compounds due to several reasons such as the volatility of some raw materials and/or the intermediate involved, since it is an open system. The key procedures in this method include repeated grinding and calcination, inert gas introduction, pelletization and so on. The purity and electrochemical behavior of cathode materials depend on the raw materials and growth parameters such as the temperature of calcination and exposure time.^157,158 The common starting materials are compounds such as oxides, hydrates, silicates/sulfates/chlorides, oxalates (M_mEt_n, salts of ethanedioic acid), acetates (M_mAc_n, salts of acetic acid), carbonates etc. which contain lithium, iron, manganese, cobalt and silicate elements. Sometimes tetraethyl orthosilicate (TEOS, Si(OC₂H₅)₄) and Li₂SiO₃ are also used. Carbon sources such as carbon black, sucrose, glucose, ascorbic acid, coal pitch,⁸¹ carbon nanospheres (CNSs), carbon nanotubes (CNTs),¹⁵⁹ reduced graphene oxide (rGO),¹²⁷ graphene oxide (GO),¹²⁵ PEG-600,¹²⁵ cellulose acetate,¹²⁵ adipic acid (C₆H₁₀O₄, etc.), acetylene black or Ketjen black^73,160 can be added to increase the electronic conductivity of the cathodes. To the best of our knowledge, there were no publications related to the synthesis of Li₂CoSiO₄ and Li₂NiSO₄ via solid-state methods (but there are via other methods). For this technique, the normal synthesis temperature for Li₂MSiO₄ ranges from 600–900 °C for 8–12 h. The initial discharge capacity ranges from 102–220.4 mA h g⁻¹ (maximum 1.33 Li⁺ extracted) at 0.06–1.0 C up to 100 cycles. The most common morphology of resultant materials is agglomerate, while a few are spheres or tubes etc., depending on the raw materials and carbon sources. The particle size varies from 7–500 nm and substitution ions include V²⁺, etc.^{65,114,161,162}

Among the cathode materials obtained via solid-state methods, Li₂Fe_0.95V_0.05SiO₄/C exhibited the highest initial discharge capacity of 220.4 mA h g⁻¹ (∼1.33 Li⁺ extracted LFS) at 0.1 C. After 50 cycles, it decreased to 146.6 mA h g⁻¹ with a retention of 78.7%. It was prepared using CH₃COOLi·2H₂O, FeC₂O₄·2H₂O, TEOS, V₂O₅ and sucrose at 650 °C for 10 h.¹⁶¹ An appropriate percentage of V-substitution (5 mol%) could possibly make for >1 Li⁺ extraction.

Nanospherical Li₂MnSiO₄/C/graphene exhibited a high initial discharge capacity of 215.3 mA h g⁻¹ (∼1.3 Li⁺ extracted LMS) at 0.05 C (Fig. 4a). After 40 cycles, this decreased to 175 mA h g⁻¹. It was prepared using MnCO₃, LiOH·H₂O, nano-SiO₂, polyethylene glycol-600 (PEG-600), glucose, cellulose acetate, and graphene oxide at 700 °C for 10 h. Spherical SiO₂, mixed carbon sources and PEG-600 form a 3D nest-like carbon network, which are favourable for improving the capacity and cyclic stability.¹²⁵

Fig. 4 (a) Charge–discharge profile of Li₂MnSiO₄/C/graphene (0.05 C) via a solid-state method;¹²⁵ (b) cyclic performance of Li₂MnSiO₄/C via a sol–gel technique at different current densities;¹³⁰ (c) charge–discharge capacity of PEDOT/Li₂MnSiO₄ via a supercritical technique at 0.05 C and at 40 °C;¹⁵⁶ (d) charge–discharge profile of Li₂MnSiO₄/C via a microwave method at C/20 and 55 °C.¹⁷¹

There are some dissatisfying results reported for the Li₂MSiO₄ (M = Mn, Co, …) series prepared via solid-state methods where >1 Li⁺ extraction (i.e. capacity > 166 mA h g⁻¹) could not be attained. Li₂MnSiO₄/C exhibited a low initial discharge capacity of 129 mA h g⁻¹ at 0.06 C, which was prepared using Li₂SiO₃, Mn(CH₃COO)₂·4H₂O and sucrose at 800 °C for 12 h.¹²⁸ Another example is Li₂MnSiO₄ prepared from LiCH₃COO·2H₂O, Mn(CH₃COO)₂·4H₂O, SiO₂ and glucose at 700 °C for 10 h. It exhibited a low initial discharge capacity of 152 mA h g⁻¹.¹⁶³ The poor performance could be ascribed to impurities in the samples, unsuitability of the carbon source and a process not effectively hampering the crystal structure destruction.

C.2. Sol–gel technique

This technique involves the initial hydrolysis of a homogeneous solution of one or more selected alkoxides (salts of adipic acid such as C₆H₅FeO₇ and SiC₈H₂₀O₄, etc.) followed by condensation. Chelating/complexation agents such as citric acid,^41,61 sucrose,^82,147 polyethylene–poly(ethylene glycol),¹⁵³ polyvinylalcohol (PVA),^148,164 tartaric acid⁸⁰ and ascorbic acid¹⁶⁵ are generally used in the sol–gel technique. The synthesis temperature usually ranges from 600–800 °C with a dwell time of 5–15 h. The morphology of Li₂MSiO₄ obtained via the sol–gel technique can be spherical,^137,164 mesoporous,^136,166 agglomerate,^61,63,167 honeycomb¹⁶⁸ etc., depending on the use of chelating agents.

Most of the mesoporous structures of Li₂MSiO₄ prepared via the sol–gel method exhibited better discharge capacities than the other structures. Hydrochloric acids and propylene oxides used for enhancing the hydrolysis of TEOS could improve the capacity retention more than the other chelating agents and catalysts.^137,138 Nanoparticles ranging from 10–200 nm could be attained by this technique with substitution with elements such as V²⁺,⁶³ Mg²⁺,⁶¹ Na⁺,¹⁶⁷ Cu²⁺,¹⁶⁹ P⁵⁺,¹⁷⁰ Zn²⁺ (ref. 169) and Ni²⁺ (ref. 169) etc.

The Li₂MnSiO₄/C composite exhibited an initial discharge capacity of 240 mA h g⁻¹ (∼1.44 Li⁺ extracted LMS) at 0.02 C (45.4% retention) and 125 mA h g⁻¹ at 0.4 C (52% retention) as shown in Fig. 4b. The discharge capacity decreased with the rate capacity, while the retention increased. It was prepared using LiCH₃COO·2H₂O, pre-synthesized Mn₃O₄ and TEOS (2.040 : 0.763 : 2.083 wt%,) as starting materials, adding acetic acid (catalyst) and sucrose (carbon source) and heating at 700 °C for 10 h. Replacing the soluble metal source with manganese oxide improved the cycle performance due to the more compact carbon-coating and more effective prevention of the side reaction between the cathode and the electrolyte.¹³⁰

The Li₂MnSiO₄/C nanocomposite exhibited a discharge capacity of 253.4 mA h g⁻¹ (∼1.53 Li⁺ extracted LMS) at 0.03 C and 149.9 mA h g⁻¹ at 1 C (56.4% retention). It was prepared using LiCH₃COO·2H₂O (0.02 mol), Mn(CH₃COO)₂·4H₂O (0.01 mol), TEOS (0.02 : 0.01 : 0.01, mol%), water–acetic acid solution (1 : 2, wt%) and glucose at 650 °C for 6 h.¹³³ The capacity decrease may be associated with the amorphous tendency causing a volumetric effect.^107,133,171

Li₂(Mn_0.94Cr_0.06)SiO₄/C nanofibers exhibited a discharge capacity of 295 mA h g⁻¹ (∼1.77 Li⁺ extracted LMS) at the first cycle and reached the highest capacity of 314 mA h g⁻¹ (∼1.88 Li⁺ extracted LMS) at the 5^th cycle. After 20 cycles, the discharge capacity was maintained at 273 mA h g⁻¹ (∼1.63 Li⁺ extracted LMS). It was prepared using N,N-dimethylformamide (DMF), Mn(CH₃COO)₂·4H₂O, LiCH₃COO·2H₂O and TEOS at 700 °C for 12 h. Appropriate Cr-substitution improved its cycling behavior due to the large unit cell volume. The carbon nanofiber matrix contributed to the faster electron and ion transportation, leading to good reversibility for the cathode materials.⁶⁴ There are also some dissatisfying results reported in the Li₂MSiO₄ (M = Mn, Co, …) series prepared via the sol–gel method where >1 Li⁺ extraction (i.e. capacity >166 mA h g⁻¹) could not be attained. Li₂MnSiO₄ exhibited a low initial discharge capacity of 113 mA h g⁻¹ at the first cycle. It was prepared using LiCH₃COO·2H₂O, Mn(CH₃COO)₂·4H₂O, TEOS and adipic acid at 700 °C.¹³⁸ Li₂CoSiO₄ exhibited only an initial discharge capacity of 32 m h g⁻¹ at 0.02 C in the 1^st cycle and it decreased after 10 cycles. It was prepared using LiNO₃, Co(NO₃)₂·6H₂O, silica particles and polyacrylic acid (PAA, used as a chelating agent) at 800 °C. The low capacity could be ascribed to poor electronic conductivity, insufficient ball milling and an unsuitable electrolyte.¹⁶⁴

C.3. Hydrothermal/solvothermal/supercritical fluid techniques

These techniques are simple, clean, of low cost and low energy consumption, being suitable to prepare high quality polycrystalline materials in closed systems, i.e. large surface area materials with the desired morphology under a relatively lower temperature/pressure than those obtained via solid-state or sol–gel techniques. The precursor concentration affects the particles’ morphology and size.^103,172,173 Nanoparticles are more active materials with regards to the electrochemical properties due to the improved electrical conductance of the electrode, shorter diffusion paths for lithium ions within a particle and increased contact area with the electrolyte.^162,174 Normally, hydrothermal/solvothermal processes are conducted in solvent at 150–200 °C for 24–96 h. For complete crystallization as well as for carbon coating, products need further annealing at 600–700 °C for 2–10 h. The supercritical fluid techniques are usually run at 300–400 °C and 30–40 MPa with a very short duration of 5–10 min. Different morphologies are observed from these techniques such as flower,^116,175 spherical,^{118,145,173,176,177} hollow sphere,¹⁷⁸ hierarchical,^120,179 hot-dog,¹⁸⁰ rod^175,177 etc., depending on the process conditions. The particle size varies from 5–400 nm. Chelating agents such as citric acid and ethylene glycol etc. were used in the hydrothermal method^{37,117,118,120,176} and some organic materials such as cetrimonium bromide, ascorbic acid, urea etc. were used in the solvothermal method.⁹⁷

By the hydrothermal method, Mali et al. first reported three polymorphs of Li₂MnSiO₄ and their NMR spectra.⁹⁷ The flower-like Li₂MnSiO₄ exhibited an initial discharge capacity of 226 mA h g⁻¹ (∼1.36 extracted LMS) at 0.05 C. After 10 cycles, the capacity was stabilized at ∼130 mA h g⁻¹. It was prepared using LiOH·2H₂O, Mn(CH₃COO)₂·4H₂O and TEOS at 700 °C for 10 h. The flower-like particles could produce good electrochemical properties due to their large surface area, unique morphology and high interconnectivity with each other.¹¹⁶

Ni-modified Li₂MnSiO₄ (LMS@Ni/C) prepared by the solvothermal technique exhibited an initial discharge capacity of 274.5 mA h g⁻¹ (∼1.67 Li⁺ extracted LMS) at 0.05 C. After 20 cycles, this decreased to 119.8 mA h g⁻¹. It was prepared using C₂H₃O₂Li·2H₂O, NiC₄H₆O₄·4H₂O, TEOS, MnC₄H₆O₄·4H₂O, cetrimonium bromide (2.2 : 0.05 : 1.0 : 0.95 : 0.1, mol%) and starch at 700 °C for 10 h. The Ni improved the ion diffusion coefficient and electronic conductivity.⁹⁷

Li₂MnSiO₄ prepared by the supercritical technique, exhibited a very high discharge capacity of 313 mA h g⁻¹ at 40 °C and 0.05 C (∼1.9 Li⁺ extracted LMS) as shown in Fig. 4c. After 20 cycles, this decreased to 240 mA h g⁻¹. It was prepared using LiOH·2H₂O, MnCl₂·2H₂O, TEOS and ascorbic acid (1 : 4 : 1 : 0.1, mol%) at 300 °C and 38 MPa for 5 min.¹⁵⁶ Nano-Li₂MSiO₄ (M = Fe, Mn, Co) materials were prepared at 300–380 °C and 38 MPa for 5–30 min. They had a high discharge capacity in the 1^st cycle at different rates but failed to maintain the stability.^145,156,177 Li₂CoSiO₄ exhibited a discharge capacity of 107 mA h g⁻¹ (∼1.51 Li⁺ extracted LCS) at 0.01 C. After 3 cycles, this decreased to 80 mA h g⁻¹. It was prepared using LiOH·2H₂O, CoCl₂·6H₂O, TEOS (4 : 1 : 1, mol%) and oleylamine at 300–350 °C and 30 MPa for 10–30 min.¹⁴⁵

There are a few results reported in the Li₂MSiO₄ (M = Mn, Co, Ni, …) series prepared via hydrothermal/solvothermal methods where >1 Li⁺ extraction (i.e. capacity > 166 mA h g⁻¹) could not be attained. Li₂Mn_0.8Ni₂SiO₄ prepared via the solvothermal technique exhibited a low discharge capacity of 100 mA h g⁻¹ (∼0.7 Li⁺ extracted LMS) at 0.2 C. It was prepared using LiOH·2H₂O, NiCl₂·6H₂O, MnCl₂·4H₂O, SiO₂ (0.386 : 0.2 : 0.8 : 0.138, wt%) and sucrose at 700 °C for 10 h. Ni-substitution maintained the structure and improved the charge capacity, but caused no improvement to the discharge capacity.¹⁸¹

C.4. Microwave method

The target materials directly absorb the electromagnetic/microwave energy (self-heating) within a short period of time, which is lower than for the other synthesis methods.¹⁸² This method is very fast, facile and high efficient. Nanoagglomerates or spheres could be prepared at 600–700 °C for 6–20 h. Cathodes with high ionic conduction were obtained above RT (60 °C).^171,183,184 By this technique, it’s difficult to extract more than 1 Li⁺ from Li₂FeSiO₄ (RT) but this was achieved with Li₂MnSiO₄.

By a microwave-assisted hydrothermal technique, Li₂MnSiO₄/C exhibited an initial discharge capacity of 210 mA h g⁻¹ (∼1.27 Li⁺ extracted LMS) at RT and 250 mA h g⁻¹ (∼1.5 Li⁺ extracted LMS) at 55 °C for 0.2 C (Fig. 4d). After 20 cycles, this drastically decreased by 50% (105 mA h g⁻¹) at RT and 15% (240 mA h g⁻¹) at 55 °C. It was prepared using LiOH·2H₂O, Mn(CH₃COO)₂·4H₂O, TEOS and sucrose at 650 °C for 6 h.¹⁷¹

Li₂MnSiO₄ obtained via a microwave-assisted solvothermal synthesis exhibited an initial discharge capacity of 260 mA h g⁻¹ at 50 °C (∼1.57 Li⁺ extracted LMS) and 130 mA h g⁻¹ at RT at 0.1 C. After 4 cycles, this drastically decreased. It was prepared using LiOH·2H₂O, Mn(CH₃COO)₂·4H₂O, TEOS and urea at 700° for 5 h.¹⁸⁵

C.5. Spray pyrolysis/combustion/hydrochemical techniques

The synthesis temperature and time vary from 600–800 °C and 4–12 h to form agglomerate/porous/spherical morphologies, depending on the conditions and carbon sources.

The spray pyrolysis method is used to prepare powders with an uniform chemical composition and a narrow particle size distribution from the nanometre to the micrometer scale, easily used for large-scale applications.¹⁸⁶ Li₂MnSiO₄/C exhibited an initial discharge capacity of 184 mA h g⁻¹ (∼1.1 Li⁺ extracted LMS) at RT and 0.05 C. After 20 cycles, this decreased to 110.4 mA h g⁻¹ (∼60%). Moreover, it exhibited an initial discharge capacity of 225 mA h g⁻¹ (∼1.36 Li⁺ extracted LMS) at 60 °C and 0.05 C but this decreased drastically at the end of the 20^th cycle.¹³⁹

Li₂MSiO₄ (M = Fe, Mn) was prepared via a combustion method, using LiNO₃, Mn(CH₃COO)₂·4H₂O, TEOS, citric acid and 20 wt% of acetylene black at 700 °C for 10 h.¹⁴² Li₂MnSiO₄/C exhibited an initial discharge capacity of 161 mA h g⁻¹ (∼0.96 Li⁺ extracted LMS) at 0.06 C.

By the hydrochemical technique, Li₂FeSiO₄/C/CNS with a double carbon coating (glucose-derived carbon and carbon nanospheres) exhibited an initial discharge capacity of 164.7 mA h g⁻¹ (∼1.0 Li⁺ extracted LFS) at 0.1 C. After 60 cycles, 98.4% capacity remained. It was prepared using CH₃COOLi·2H₂O, Fe(CH₃COO)₂·4H₂O, TEOS, glucose and CNS at 700 °C for 12 h. The double carbon coating increased electron transport among the particles.¹⁸⁷

D. First principles density functional theory (DFT) calculations

First principles density functional theory (DFT) is an approach to the quantum mechanical many-body problem, where the system of interacting electrons is mapped onto an effective non-interacting system with the same total density. The total energies of all the compounds were usually calculated using the Projector Augmented Wave (PAW) formalism as implemented in the Vienna Ab initio Simulation Package (VASP), or the full-potential linear augmented plane wave (FP-LAPW) method, with the exchange and correlation energies approximated in the Perdew–Burke–Ernzerh generalized gradient approximation (PBE-GGA), local spin density approximation (LSDA), generalized gradient and local density approximations plus an on-site Coulomb self-interaction correction potential/ultrasoft pseudopotential (GGA + U^SIC and LDA + U^SIC, respectively). Two versions of spin-polarized (ferromagnetic/anti-ferromagnetic, i.e. FM/AFM) calculations were employed.^{44,96,201–205}

Through DFT calculations for a specific structure, the properties of Li₂MSiO₄ (M = Mn, Fe, Co, Ni, …) can be accurately predicted and its reaction mechanisms can be fully described. The applications of DFT calculations include crystal structure modeling and stability investigations of delithiated and lithiated phases, averaging of the lithium intercalation voltage, prediction of charge distributions and band structures, and kinetic studies of lithium ion diffusion processes (Li⁺ vacancy migration barriers^86,206), which can provide atomic understanding of the capacity, reaction mechanism, rate capacity, and cycling ability. The results obtained from DFT are valuable for establishing the relationship between the structure and the electrochemical properties, promoting the design of new electrode materials.²⁰¹

The total energy vs. volume of Li₂MnSiO₄ polymorphs (orthorhombic β_II-Pmn2₁/γ_II-Pmnb and monoclinic γ_s-P2₁/n) and their electrochemical properties as electrodes for LIBs were investigated by Arroyo-de Dompablo et al.,⁴⁴ combining experimental and computational methods. Calculation results show that the crystal structure has little effect on the average Li⁺ intercalation voltage (4.18/4.19/4.08 V for Pmnb/Pmn2₁/P21/n, respectively). The Pmnb form is 2.4/65 meV/f.u. more stable than the Pmn2₁/P2₁/n forms, respectively (Fig. 5, GGA + U results) but one cannot deny that other variants of β- and/or γ-Li₃PO₄ could be energetically accessible. In fact, a β-Li₃PO₄ (β_I-Pbn2₁) type Li₂MnSiO₄ sample had been prepared through the hydrothermal method. DFT calculations also revealed that the denser Pmn2₁ polymorph can be obtained by a high-pressure high-temperature treatment of the other polymorphs or their mixtures.⁴⁴

Fig. 5 Calculated total energy vs. volume curves of Li₂MnSiO₄ polymorphs; Pmn2₁ (red), Pmnb (blue) and P2₁/n (green). The calculated average voltage for the 2 electron process is given in parentheses.^44,209

⁶Li MAS NMR spectroscopy combined with DFT calculations was used to study the structural differences between various polymorphs of Li₂MSiO₄ (M = Mn, Fe, Zn). Results show that fully lithiated Li₂MSiO₄, delithiated LiMSiO₄ and MSiO₄ are semiconducting and the band gap of Li₂MSiO₄ decreases while extracting lithium ions.^96,97 The Si–O bonds remain almost unchanged during lithiation–delithiation for all polymorphisms, which contributes significantly to the structural stability.⁸⁶ Comparisons exhibit a more promising role for the monoclinic P2₁/n configuration. The corresponding fully delithiated MSiO₄ attained a better stability due to the high-spin state from M²⁺ to M³⁺ and to M⁴⁺ ions.⁴⁸ In contrast to the LiMPO₄ counterpart, the potentials of Li₂MSiO₄ are largely increased due to the higher valence state of the M³⁺/M⁴⁺ redox couple²⁰⁷ and Li⁺ conduction exhibits a two-dimensional anisotropic character.¹⁹

Kokalj et al. predicted through DFT calculations that it might be possible to obtain a stable material with >1 Li⁺/f.u. by using a Li₂Mn_xFe_1−xSiO₄ solid solution. The voltage required is lower by 0.7 V than that for the pure Fe counterpart.²⁰² However, Larsson et al. argued that when the ratio of Mn substitution was lowered to 12.5%, the structural distortion and high voltage would destroy the feasibility of this design.²⁰⁸

Kalantarian et al. showed by DFT calculations that β_II-Li₂MnSiO₄ (Pmn2₁) should have better electrochemical properties, even after a 2 Li⁺/f.u. extraction.²⁰³ The insertion–extraction mechanism was put into a relation with the voltage–capacity behaviour by considering a diffusion model. They also related the voltage behaviour to relevant parameters such as the reaction energy, Li⁺ diffusion coefficient and particle size, so suggesting some strategies for optimizing materials.²⁰⁴ Li₂M_0.5N_0.5SiO₄ (M, N = Mn, Fe, Co, Ni) compounds with a Pmn2₁ structure were studied by DFT using GGA (+U^SIC) and LSDA (+U^SIC) methods. Mixed compounds score better than pure materials, except for Fe–Mn in comparison to Fe. Considering both the electron conductivity and theoretical lithiation–delithiation reaction voltage, the best properties are shown by Fe–Ni, Mn–Ni, Fe–Co and Mn–Co, in descending order. Fe–Ni is theoretically the most promising material. Based on the gravimetric specific energy, these materials are sorted as: Mn–Fe, Fe, Mn, Mn–Ni, Fe–Co, Mn–Co, Fe–Ni, Co, Co–Ni and Ni, in ascending order (Fig. 6).²⁰⁵

Fig. 6 Band gaps for delithiated states of Li₂M_0.5N_0.5SiO₄ (M, N = Mn, Fe, Co, Ni) calculated via GGA/LSDA + U^SIC methods.²⁰⁵

E. Techniques for improving the electrochemical performance

As mentioned earlier, just like LiFePO₄, Li₂MSiO₄ (M = Fe, Mn, Co, Ni), cathode materials suffer from poor electronic conductivity and a slow lithium ion diffusion rate, which inhibit their use in higher power applications.^45,210 To overcome these drawbacks, the following strategies were adopted: (a) carbon coating; (b) optimization of the particle size and morphology; and (c) ion-substitution.

E.1. Carbon coating

In synthesizing Li₂MSiO₄, carbon sources are often used to improve the electrochemical properties and at the same time act as buffers preventing grain growth and the formation of hard agglomerates during heat treatment.¹⁴² Generally, it works in two ways: (i) in situ carbon coating: the starting precursors and carbon source are mixed and then calcinated at high temperature. This type of coating would form a thin and homogeneous carbon layer on the surface of the particles and improve the electrochemical properties.²¹¹ The formed carbon can also affect the growth of particles or possibly increase impurities in the target materials. (ii) Ex situ carbon coating: carbon-free materials are prepared and then mixed with the carbon source by ball milling/grinding. Heat treatment is necessary if a carbon precursor rather than carbon black is used. This type of coating has little effect on the morphology of the particles and the amount of impurities.¹²⁸

Many researchers have suggested that the selected carbon precursors are directly involved in the characteristics of the carbon additives, in terms of their structure, distribution and the thickness of the coating layer, which are proportional to the performance of electrodes.^212–215

Carbon coating not only provides conductive connections among the active particles which are favourable for electron and ion-transfer (Fig. 7a),²¹⁶ but also decreases the particle size of the active materials effectively during the heat treatment. It can shorten the diffusion path of the lithium ions and also facilitate good contact between neighbouring particles with a reduction of the polarization.^45,61,217 Uniform carbon coating (∼2–5 nm thickness on nanoparticles) increases the capacity and rate performance by providing short pathways and huge electrochemical interfaces for fast ion diffusion and transportation (Fig. 7b).^133,218,219 Various carbon precursors directly affected the electrical conductivity of the carbon, which supports the improvement of the electrochemical performance of Li₂MSiO₄/C cathode materials.^136,220 Increased sp²-coordinated carbon promised a better electrochemical performance.^221–223 Functional structures or ring-forming structures of the organic precursor have attracted considerable attention as carbon sources for electrodes.^212,221,222

Fig. 7 (a) Schematic presentation of electron and ion transport in a carbon-coated cathode material;²¹⁶ (b) uniform thickness of carbon coating on a Li₂MSiO₄ particle.¹³³

If the aggregated carbon content is too high, it will block the pathway for electrolyte percolation. Hence the ionic conduction is blocked.²²⁴ The total amount of impurities increases with the carbon content, i.e. the content of electrochemically active Li₂MSiO₄ decreases with increasing carbon content.

Up to now, many carbon sources have been used in synthesizing Li₂MSiO₄ cathode materials, such as carbon nanospheres (CNSs),¹⁸⁷ carbon nanotubes (CNTs),¹⁵⁹ reduced graphene oxide (rGO),¹²⁷ graphene oxide (GO),¹²⁵ PEG-600,¹²⁵ cellulose acetate,¹²⁵ adipic acid,²²⁵ acetylene black or Ketjen black,^73,160 sucrose,^{20,114,128,130,131,136,139,158,170,171,180,188,212,226–229} cellulose acetate,¹⁵⁸ acetylene black,^{73,109,118,132,200,226,229–232} cellulose,¹⁰⁸ polyethylene–poly(ethylene glycol),^{125,141,153,233} ascorbic acid,^56,156 citric acid,^{41,56,63,82,132,141,191,216,227,233–236} corn starch,^84,166 glucose,^{62,81,118,125,183,187,212,237} carbon nanosheets/tubes,^{187,238–241} graphene,^125,127 pyromellitic acid,^135,242 P123 polymer,²⁴³ carbon black,^{24,106,113,134,147,148,153,187,188,244,245} polyacrylonitrile,¹⁸⁹ tartaric acid,⁸⁰ adipic acid,^{137,138,160,162,246} pitch^81,122 and phenolic resin.²¹¹

Li₂MnSiO₄ obtained via a sol–gel technique using 2.1 wt% sucrose showed a higher cyclic performance than that using 5.9 wt% content.²³⁵ Li₂MnSiO₄ obtained via a sol–gel technique using 0.2 mol% adipic acid (acts as a carbon source and chelating agent) exhibited a better discharge capacity of 113 mA h g⁻¹ than that using 0.2 mol% content (58 mA h g⁻¹).¹³⁸ Carbon suppression to crystallites and connection with the pores in Li₂MnSiO₄/C resulted in a carbon-coated compound. This showed a higher discharge capacity of 275.2 mA h g⁻¹ than that in uncoated Li₂MnSiO₄ (178.3 mA h g⁻¹).¹³⁶

Li₂MnSiO₄/C obtained via a combustion method using 0 wt%, 10 wt%, and 20 wt% of acetylene black exhibited the discharge capacity of 2 mA h g⁻¹, 128 mA h g⁻¹ and 164 mA h g⁻¹, respectively.¹⁴²

Graphitized carbon sources afford a better electrochemical performance than the non-graphitized forms. A Li₂MnSiO₄/C/graphene composite was prepared using carbon sources like polyethylene glycol 600, glucose, cellulose acetate and graphene oxide. It exhibited an initial discharge capacity of 215.3 mA h g⁻¹ at 0.05 C. The 3D nest-like carbon network and carbon layer are favourable for improving the capacity and cycling stability.¹²⁵

E.2. Optimization of the particle size and morphology

As per the diffusion formula t = L²/2D (where t is the diffusion time, L is the diffusion distance and D is the diffusion coefficient), the rate of the electrochemical reaction increases with the reduction of the particle size. The porous structure significantly shortens the diffusion time of Li⁺ in Li₂MSiO₄ cathode materials.^247,248 Nanospherical particles increase the trap density of the cathodes, because the larger size of the secondary particles provides a denser packing, while the smaller size of the primary particles improves the Li⁺ and electron conduction.¹⁷¹

Li₂MnSiO₄ particles exhibited a uniform spherical shape and a size of around 50 nm coated with a carbon/graphene layer (Fig. 8a).¹²⁵ Nanospherical Li₂MnSiO₄ (20–30 nm) was synthesized using a chelating agent (adipic acid). It exhibited an initial discharge capacity of ∼161 mA h g⁻¹ and was stable up to 50 cycles.¹³⁷

Fig. 8 (a) TEM of Li₂MnSiO₄ particles exhibiting an uniform spherical shape and a size of around 50 nm coated with a carbon/graphene layer;¹²⁵ (b) SEM of Li₂MnSiO₄ thin flowers with a flake-type petal;¹¹⁶ (c) TEM of Li₂MnSiO₄ particles with a hierarchical structure.²⁵⁰

The 20 nm thin flower-like morphology with a flake margin of petals (Fig. 8b) was formed by Li₂MnSiO₄ during the nucleation growth stage in the presence of water under an optimized hydrothermal treatment. The resulting material exhibited an electrochemical stability of ∼125 mA h g⁻¹ due to its large surface area, unique morphology and high interconnectivity beween the nearest neighbouring particles.¹¹⁶

Large numbers of pores in the mesoporous and macroporous structures (also in the hierarchical porous structure) greatly influence the cycling behavior, resulting in a discharge capacity of 275.2 mA h g⁻¹ in Li₂MnSiO₄/C and 163 mA h g⁻¹ in Li₂FeSiO₄/C at 0.06 C. The current penetrating electrolyte through the porous structure was of benefit for promoting Li⁺ transportation.^{133,136,166,178,249} Honeycomb-like Li_1.8MnSO₄ with mesopores of 6–8 nm was produced by decomposing raw materials at high temperature (to release CO₂ and H₂O), and it exhibited a discharge capacity of 110.9 mA h g⁻¹.¹⁶⁸

Li₂MnSiO₄/C/V₂O₅ in which V₂O₅ nanowires adhered to Li₂MnSiO₄/C nanoparticles exhibited a high discharge capacity of 277 mA h g⁻¹ (1.66 Li⁺ extracted LMS) at 0.1 C, because V₂O₅ nanowires suppressed the dissolution of manganese and protected Li₂MnSiO₄ from a corrosive reaction with the electrolyte.¹⁹⁵

Li₂MnSiO₄ with a hierarchical structure (200–400 nm, Fig. 8c) was prepared via a supercritical fluid process using oleic acid as a surfactant. It exhibited an initial discharge capacity of 283 mA h g⁻¹ (1.7 Li⁺ extracted LMS). The suitable particle size and the special structure enabled nearly 2 Li⁺ to be extracted.

E.3. Cation-substitution

Aliovalent/isovalent cation substitution is often employed to improve the electrochemical performance of polyanionic compounds by the alteration of the cell/particle sizes of the cathode materials.^169,251 Aliovalent ion substitution is more effective^252–254 and can induce defects (vacancies or interstitials) in the lattice by charge compensation to enhance the conductivity.^{19,56,169,252–255} A high atomic percentage of aliovalent (Al³⁺, Zr⁴⁺ and Nb⁵⁺) substitution was reported to reduce the lithium miscibility gap and increase the phase transformation kinetics^256,257 but a theoretical study on energetic grounds proved that polyanions are not tolerant to aliovalent substitution on either Li⁺ or M²⁺ sites.²⁵⁸ The effect of doping of aliovalent cations is still under debate.^{251,258–261}

The substitution of Mn²⁺ by Cr³⁺ in Li₂(Mn_1−xCr_x)SiO₄/carbon nanofiber composites (x = 0.03, 0.06 and 0.1) increased the crystal unit cell volumes. The crystalline structure was prevented from collapsing and the structural stability was improved during the charge–discharge cycles. Li₂(Mn_0.94Cr_0.06)SiO₄ exhibited a discharge capacity of 314 mA h g⁻¹ and an improved cycling behavior due to the maximized unit cell volume.⁶⁴

The substitution of Mn²⁺ by Ti⁴⁺ in Li₂(Mn_1−xTi_x)SiO₄ (x = 0, 0.06, 0.1 and 0.2) led to a slight shrinkage of the volume due to the smaller ionic radius of Ti⁴⁺ compared to that of Mn²⁺. The particles had smaller sizes, better monodispersion and larger specific areas compared to pristine Li₂MnSiO₄. Li₂(Mn_0.8Ti_0.2)SiO₄ showed a good cycling performance and maintained a capacity of around 100 mA h g⁻¹ after 50 cycles (Fig. 9a).¹⁹⁶

Fig. 9 (a) Cyclic performance of Li₂Mn_1−xTi_xSiO₄ (x = 0, 0.06, 0.1 and 0.2);²⁶² (b) cyclic performance of Mg²⁺, Ga³⁺ and Al³⁺ substituted Li₂MnSiO₄.²⁶³

After substitution of trivalent ions such as Ga³⁺, Al³⁺ and Mg²⁺ in Li₂MnSiO₄, small intensity peaks were found in the region of Li/Mn sites but the substitution did not affect the structural arrangement. Only Ga³⁺ led to a better charge–discharge capacity (Fig. 9b) and exhibited an effect on well-dispersed nanoparticle formation.⁵⁶ If substitution was on the Si site, charge compensation could be achieved through the creation of Li⁺ interstitials.¹⁹

Li₂FeSiO₄ suffers low electronic conductivity and slow Li⁺ diffusion. The substitution of iron with other transition metals is a possible solution to the problem. At the same time, the initial structure of Li₂MnSiO₄ is progressively amorphized during delithiation, which leads to poor stability during cycles. In order to combine the high discharge capacity of Li₂MnSiO₄ with the high cycling stability of Li₂FeSiO₄, partial substitution of Fe²⁺ with Mn²⁺ was employed.^47,202,233 In this way, substitution of Fe²⁺ with Mn²⁺, which undergoes a 2⁺ → 4⁺ transition, can facilitate a “>1 electron redox reaction” through the following reaction with the removal of up to 2 Li⁺, resulting in a capacity increase from 170 mA h g⁻¹ to 340 mA h g⁻¹.²⁶⁴

Li⁺₂(Fe²⁺_1−xMn²⁺_x)SiO₄ ↔ Li⁺_1−x(Fe³⁺_1−xMn⁴⁺_x)SiO₄ + (1 + x)Li⁺ + (1 + x)e⁻

The substitution of Fe²⁺ with Mn²⁺ in Li₂(Fe_1−xMn_x)SiO₄ (x = 0, 0.03, 0.1, 0.2 and 0.3) through a solution route led Li₂(Fe_0.9Mn_0.1)Si0₄/C to exhibit an initial discharge capacity of 158.1 mA h g⁻¹ (∼0.95 Li⁺ extracted LFS) with a capacity retention of 94.3% after 30 cycles.¹¹⁴

The substitution of Mn²⁺ with Fe²⁺ in Li₂(Mn_1−xFe_x)SiO₄/C (0 ≤ x ≤ 0.8) via a spray pyrolysis/ball milling/annealing route led Li₂(Mn_0.5Fe_0.5)SiO₄/C to exhibit an initial discharge capacity of 149 mA h g⁻¹ (∼0.9 Li⁺ extracted LFS) at 1 C.¹⁰⁹ In contrast, the substitution of Fe²⁺ with Mn²⁺ in Li₂(Fe_1−xMn_x)SiO₄ (x = 0, 0.3, 0.5, 0.7, 1) led to a discharge capacity with only 61.4% retention after 50 cycles.¹¹³ The Mn-substituted materials exhibited a higher redox potential and higher initial discharge capacity but suffered a poor cycling performance and electrochemical reversibility. This might be due to the structural stability and electronic conductivity. Li₂(Fe_0.8Mn_0.2)SiO₄ showed good reversibility and exhibited an initial discharge capacity of 230.1 mA h g⁻¹ (∼1.4 Li⁺ extracted LFS) at 0.1 C and retained 162 mA h g⁻¹ after 20 cycles.

Other efforts involve the replacement of O with N,²⁶⁵ substituting polyanionic SiO₄ with AsO₄,²⁶⁶ BO₃,²⁶⁷ VO₄,²⁶⁸ and doping trivalent Al and Ga on Si sites.¹⁹

F. Conclusions

The pursuit of effective cathode materials for LIBs is a challenge for current/future energy storage requirements. The lithium transition-metal orthosilicate Li₂MSiO₄ (M = Mn, Fe, Co, Ni, …) should be paid more attention because it has a high theoretical capacity (∼330 mA h g⁻¹) with a possible 2 Li⁺ extraction per formula unit. In the present perspective, an attempt has been made to review Li₂MSiO₄ materials based on the polymorphism of their structures, synthetic methods and techniques for improving their electrochemical performance.

The stable Li₂MSiO₄ polymorphs include high temperature monoclinic γ_o/γ_s (P2₁/n), high temperature orthorhombic γ_II (Pmna, Pmnb, Cmma) and low temperature orthorhombic β_I (Pbn2₁, Pna2₁) and β_II (Pmn2₁), which are influenced by the synthetic conditions and electrochemical behavior. Metastable monoclinic P_n-Li₂MnSiO₄ would convert to the β_II structure (Pmn2₁) above 370 °C. The densest β_II (Pmn2₁) polymorph attained the best electrochemical performance due to its special structure. It can be prepared from other less dense polymorphs, by controlling the temperature, pressure and/or electrochemical process.

Once the materials’ amorphization or collapse were inhibited during lithium ion insertion/extraction from the host lattice, reversible structural changes and good cycling performances could be attained. Hence, it is suggested that ion-substitution, for stabilizing the structures, is an effective way to achieve high-capacity Li₂MSiO₄ materials. Carbon-coating and optimization of the particle size/morphology are also applied as exterior modifying methods to enhance the electrochemical performance. Thus understanding the relationship between the polymorph structure, synthesis and phase transition in Li₂MSiO₄ is most important.

Acknowledgements

The authors gratefully thank Mr Z.-S. Gao, Mr D. Chen, Mr J. Li, Ms G.-G. Zhao, Mr J.-H. Liu and Ms H.-F. Zhang in our group. This work was supported by Postdoctoral Research Award Foundation of Wuhan University of Technology (WUT), and Open Research Foundation of State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, WUT, China (Grant Nos. 2014-KF-6 & 2015-KF-4).

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