WO2013180658A1 - Glass ceramic electrolyte system - Google Patents

Abstract

The invention relates to a glass ceramic electrolyte system, and in particular, to a doped lithium- aluminium-germanium-phosphate (Li₂O-Al₂O₃-GeO₂-P₂O₅) electrolyte. A method of preparing the glass ceramic electrolyte system and its use in lithium-ion battery are also disclosed.

Description

GLASS CERAMIC ELECTROLYTE SYSTEM Cross-Reference to Related Application

[0001] This application claims the benefit of priority of United States of America Provisional Patent Application No. 61/654,363, filed June 1, 2012, the contents of which being hereby incorporated by reference in its entirety for all purposes.

Technical Field

[0002] The invention relates to a glass ceramic electrolyte system, and in particular, to a doped lithium-aluminium-germanium-phosphate (Li₂0-Al₂0₃-Ge0₂-P₂0₅) electrolyte. A method of preparing the glass ceramic electrolyte system and its use in lithium-ion battery are also disclosed.

Background

[0003] The state-of-art lithium-ion batteries have advanced applications in automotive and large scale energy storage systems. However, such batteries using organic liquid electrolytes are plagued with issues associated with limited performance, thermal and ambient environment instability, flammability and corrosiveness with the lithium metal. Lithium-air batteries with aqueous cathodes also need electrolytes that are stable against Li and H₂0, with conductivity a > 10^A S cm^"1. Since 1970s, solid-state electrolytes (SSEs) have been the subject of intense research replacing liquid electrolytes in lithium-ion/air batteries, which anticipated improvements in performance, safety, enhanced geometric flexibility (i.e. size/shape), and ease of manufacture (e.g. bipolar configurations). SSEs, in the forms of metallate, phosphate, sulphate, silicate and borate, satisfy numerous requirements, including: superionic (fast ionic) transport, negligible electronic conduction and thermodynamic stability over a wide range of temperatures.

[0004] Among the Li analogue of a NASICON-type structure, Li_xGe₂(P0₄)₃ (0 <x <3) has been reported to exhibit corner-sharing Ge0₆ octahedra (i.e. Ge replaced by any tri-, tetra- and penta- valent ion) and P0₄ tetrahedra, leading to a three-dimensional (3D) framework of Ge₂(P0₄)₃ formulae with interconnected conducting channels where lithium cations occupy two different sites usually noted Ml at 6b(0, 0, 0) and M2 at 18e(x, 0, 1/4) wickoff positions. The interconnected voids of the NaSICON-structure allow the fast ionic transport, which depends upon the size of framework, activation energy required for the movement of Li⁺ ions and lattice parameters, whose value can be modified by changing the compositions. The transition from normal conducting to high conducting behaviour involves disordering of one of the ion sub-lattices of the material. At low temperatures, all the ions are situated on well-defined lattice sites and have very low mobility.

As the temperature is increased, the mobile ions begin populating the interstitial sites. In the superionic phase, these ions are distributed over a large number of available sites. NaSICON phase of superionic Li_1+xAl_xGe₂._x(P0₄)₃ (LAGP) glass ceramic electrolytes (GCEs) are desirable because of their high ionic conductivity (~ 4.62 x 10^~3 S cm^"1 at 27 °C) and stability against Li metal. Partial substitution of Ge⁴⁺ with Al³⁺ has shown improved density of LiGe₂(P04)3 as well as enhanced Li⁺ ion conductivity.

[0005] To enhance the conductivity further in NaSICON-structure, the following approaches can be adopted: (i) optimization of preparative parameters, (ii) stabilization of 3D-tunnel structure, through which Li⁺ ions can migrate, (iii) increasing the disorder of the system by introducing dielectric phase impurity in the parent glass. This has stimulated interest in investigating a new strategy in material research introducing high ionic conductivity and offers great potential in rechargeable lithium-ion/air batteries.

Summary

[0006] The inventors have surprising found out that by introducing a glass former in a lithium- aluminium-germanium-phosphate (LAGP) electrolyte system, the amorphous glass structure is stabilized. As a result, the ionic conductivity of the doped LAGP electrolyte system is about three times higher than that of the bare LAGP electrolyte system, which is attributed to a higher defect density and Li ion mobility.

[0007] Thus, in a first aspect, there is provided a glass ceramic electrolyte system, comprising a lithium-aluminium-germanium-phosphate (Li₂0-Al₂0₃-Ge0₂-P₂05) electrolyte doped with boron oxide (B2O3). In various embodiments, the glass ceramic electrolyte system contains at most about

0.5 wt% B₂0₃, based on the total weight of the electrolyte system, such as about 0.3 wt% B₂0₃. [0008] In a second aspect, a method for preparing a glass ceramic electrolyte system comprising a lithium-aluminium-germanium-phosphate (Li₂0-Al₂0₃-Ge0₂-P₂05) electrolyte doped with boron oxide (B₂O₃) is disclosed. The method comprises:

milling a mixture of lithium precursor, aluminium precursor, germanium precursor, and phosphate precursor in a solvent to obtain a slurry;

drying the slurry to obtain a powder mixture;

heating the powder mixture until the powder mixture is melted;

quenching the melted powder mixture;

annealing the quenched powder mixture to obtain a lithium-aluminium-germanium- phosphate (Li₂0-Al₂0₃-Ge0₂-P₂0₅) glass;

doping the Li₂0-Al₂0₃-Ge0₂-P₂0₅ glass with Β₂θ3 or a boron oxide precursor to obtain the Li₂0-Al₂0₃-Ge0₂-P₂0₅ electrolyte system.

[0009] In a third aspect, a lithium-ion battery comprising a glass ceramic electrolyte system sandwiched between an anode and a cathode, wherein the glass ceramic electrolyte system comprises a lithium-aluminium-germanium-phosphate (Li₂0-Al₂0₃-Ge0₂-P₂0₅) electrolyte doped with boron oxide (B₂0₃) is disclosed.

Brief Description of the Drawings [0010] In the drawings, like reference characters generally refer to the same parts throughout the different views. The drawings are not necessarily drawn to scale, emphasis instead generally being placed upon illustrating the principles of various embodiments. In the following description, various embodiments of the invention are described with reference to the following drawings.

[0011] Fig. 1 shows DSC for the LAGP-xB₂0₃ (x = 0.0-0.4 wt.%) glass, indicating crystallization and kinetic behaviour.

[0012] Fig. 2 shows X-ray diffraction patterns of LAGP GCEs sintered at various temperatures from 825-875 °C for 12 h.

[0013] Fig. 3 shows X-ray diffraction patterns of LAGP-xB₂0₃ (x = 0.0-0.4 wt. %) GCEs sintered at 850 °C for 12 h.

[0014] Fig. 4 shows FESEM images of fractured surface of LAGP GCEs sintered at various temperatures (a) 825 °C, (b) 850 °C, and (c) 875 °C for 12 h.

[0015] Fig. 5 shows FESEM images of fractured surface of LAGP-xB₂0₃ (x = 0.0-0.4 wt.%) GCEs sintered at 850 °C for 12 h: (a) 0.1 wt.%, (b) 0.2 wt.%, (c) 0.3 wt.%, and (d) 0.4 wt.%.

[0016] Fig. 6 shows static ⁷Li nuclear magnetic resonance (NMR) spectra of rhombohedral LAGP-xB₂0₃ (x = 0.0-0.4 wt.%) phases at room temperature.

[0017] Fig. 7 shows room temperature impedance spectra of LAGP GCEs sintered at various temperatures from 825-875 °C for 12 h. [0018] Fig. 8 shows room temperature impedance spectra of LAGP-xB₂0₃ (x = 0.0-0.4 wt.%) GCEs sintered at temperatures 850 °C for 12 h.

[0019] Fig. 9 shows the Arrhenius plots of conductivity of parent LAGP GCEs sintered at temperature 825-875 °C for 12 h in the temperature range -30 to 100 °C.

[0020] Fig. 10 shows the Arrhenius plots of conductivity of LAGP-xB₂0₃ (x = 0.0-0.4 wt.%) GCEs sintered at 850 °C for 12 h in the temperature range -30 to 100 °C.

[0021] Fig. 11 shows variation of interfacial resistance of LAGP-xB₂0₃ (x = 0.3 wt.%) GCEs with storage time (Li/GCE/Li cells, frequency range 10 mHz to 1 MHz, amplitude 100 mV).

Description

[0022] The following detailed description refers to the accompanying drawings that show, by way of illustration, specific details and embodiments in which the invention may be practised. These embodiments are described in sufficient detail to enable those skilled in the art to practise the invention. Other embodiments may be utilized and structural, logical, and electrical changes may be made without departing from the scope of the invention. The various embodiments are not necessarily mutually exclusive, as some embodiments can be combined with one or more other embodiments to form new embodiments.

[0023] Water-stable superionic lithium-aluminium-germanium-phosphate (Li₂0-Al₂0₃-Ge0₂-

P₂0₅) (LAGP) glass ceramic electrolyte (GCE) system is successfully prepared by melting and quenching method. A novel vitreous phase is prepared by introducing boron oxide (B₂0₃) as a glass former in the bare parent LAGP system to stabilize the amorphous glass structure.

[0024] Thus, in accordance with a first aspect of the invention, a glass ceramic electrolyte system comprising a LAGP electrolyte doped with B₂0₃ is provided.

[0025] In the present context, the expression "LAGP-xB₂0₃" is used to designate the present LAGP GCE system where "x" denotes the relative weight amount of the dopant B₂0₃ present in the system, expressed as a weight percent (wt.%) based on the total weight of the electrolyte system.

[0026] In various embodiments, the glass ceramic electrolyte system contains at most about 0.5 wt% B₂0₃, based on the total weight of the electrolyte system. For example, the glass ceramic electrolyte system contains about 0.5 wt% B₂0₃ (i.e. LAGP-0.5B₂O₃), about 0.4 wt% B₂0₃ (i.e. LAGP-0.4B₂O₃), about 0.3 wt% B₂0₃ (i.e. LAGP-0.3B₂O₃), about 0.2 wt% B₂0₃ (i.e. LAGP- 0.2B₂O₃), about 0.1 wt% B₂0₃ (i.e. LAGP-0.1B₂O₃), or less.

[0027] In one embodiment, the glass ceramic electrolyte system contains about 0.3 wt% B₂0₃ (i.e. LAGP-0.3B₂O₃).

[0028] To verify the lithium ion conductivity, stability, morphology and other properties of the present LAGP GCE systems, the so-prepared LAGP-xB₂0₃ (x = 0-0.4 wt.%) electrolytes are systematically investigated by differential scanning calorimetry (DSC), X-ray diffraction (XRD), field-emission scanning electron microscope (FE-SEM), nuclear magnetic resonance (NMR), and ionic conductivity (σ) studies in a wide temperature range (-30-100 °C).

[0029] Highest conductivity of ~ 3.2 x 10^"3 S cm^"1 at 30 °C was obtained for bare parent LAGP GCEs sintered at 850 °C for 12 h. An optimum amount of B₂0₃ (x = 0.3) addition to the LAGP system shows very high structural as well as thermal stability and a ~ 8.75 x 10^"3 S cm^"1 at 30 °C, which is about three times higher than that of the bare parent LAGP, which is attributed to higher defect density and Li ion mobility. In most cases, two regions of activation energy (0.58 - 0.74 eV) and (0.18 - 0.42 eV) for lithium transport are observed. It can be inferred that B₂0₃-added LAGP electrolytes in appropriate compositions would be promising candidates for an all-solid-state battery because of its facile preparation, relatively high Li ion conductivity, and good interfacial stability against Li metal.

[0030] The bare parent LAGP GCEs system was prepared by a melting and quenching method, and green pellets were sintered at various sintering temperatures to optimize the highest ionic conductivity for the GCEs.

[0031] Thus, in accordance with a second aspect of the invention, a method for preparing a glass ceramic electrolyte system comprising a LAGP electrolyte doped with B₂0₃ is provided. The method comprises: milling a mixture of lithium precursor, aluminium precursor, germanium precursor, and phosphate precursor in a solvent to obtain a slurry;

drying the slurry to obtain a powder mixture;

heating the powder mixture until the powder mixture is melted;

quenching the melted powder mixture;

annealing the quenched powder mixture to obtain a lithium-aluminium-germanium- phosphate (Li₂0-Al₂03-Ge0₂-P₂0₅) glass;

doping the Li₂0-Al₂0₃-Ge0₂-P₂C>5 glass with B₂0₃ or a boron oxide precursor to obtain the Li₂0-Al₂0₃-Ge0₂-P₂05 electrolyte system.

[0032] In the milling step, the lithium precursor, aluminium precursor, germanium precursor, and phosphate precursor are ground into smaller pieces in a milling machine, such as but is not limited to a horizontal planetary mill. The mixture is milled in a solvent such as an alcohol for a period of time, for example 1 h, 3h, 6 h, 12 h, 18h, 24h, 30 h, 36 h, or more, to obtain a slurry. In various embodiments, the solid content in the milling mixture is about 10 vol.%, based on the total volume of the milling mixture.

[0033] The respective precursors are added in stoichiometric ratio in accordance to the LAGP formula Li] .5Al₀.5Gei.5(PO₄)₃. However, due to the volative nature of lithium compounds at high melting temperature and in order to maintain the stoichiometry of the LAGP formula, an extra amount of lithium precursor has been added to the milling mixture. In various embodiments, an extra 1-5 wt% of lithium precursor may be added to the milling mixture, such as, 1 wt%, 2 wt%, 3 wt%, 4 wt%, or 5 wt%.

[0034] In various embodiments, the lithium precursor is selected from the group consisting of Li₂C0₃, Li₂0, LiOH, Li(CH₃COO), Li₂C₂0₄, LiN0₃, LiH₂P0₄, Li₂HP0₄, Li₃P0₄, Li₂B₄0₇ and a combination thereof.

[0035] In one embodiment, the lithium precursor is Li₂C0₃.

[0036] In various embodiments, the aluminium precursor is selected from the group consisting of A1₂0₃ , A1(N0₃)₃, Al(OH)₃, AIOOH, A1₂(C₂0₄)₃, Al(CH₃COO)₃, A1₂(B₄0₇)₃ and a combination thereof.

[0037] In one embodiment, the aluminium precursor is A1₂0₃.

[0038] In various embodiments, the germanium precursor is selected from the group consisting of Ge0₂, GeO, GeO(OH)₂, Ge(N0₃)₂, Ge(N0₃)₄, GeC₂0₄, Ge(C₂0₄)₂, Ge(CH₃COO)₂,

Ge(CH₃COO)₄, GeB₄0₇ and a combination thereof.

[0039] In one embodiment, the germanium precursor is Ge0₂.

[0040] In various embodiments, the phosphate precursor is selected from the group consisting of NH₄H₂P0₄, P₂0₅, P₄0₆, P₄0₇, P₄0₈, P₄0₉, PO, P₂0₆, P₂0₃, (NH₄)₃P0₄, (NH₄)₂(HP0₄) and a combination thereof. [0041] In one embodiment, the phosphate precursor is NH₄H₂P0₄.

[0042] In various embodiments, the boron oxide precursor is selected from the group consisting of B(N0₃)₃, B(CH₃COO)₃, B₂(C₂0₄)₃, (NH₄)₂(B₄0₇), NH₄(HB₄0₇), B(H₂P0₄)₃, B₂(HP0₄)₃, BP0₄, B(A10₂)₃, B(GeO(OH)₃)₃, B₂(Ge0₂(OH)₂)₃, B(Ge(OH)₇), B(Ge(OH)₃)₃, B₂(GeO(OH)₂)₃,

B(Ge(OH)₅) and a combination thereof.

[0043] After milling, the slurry is dried to obtain a powder mixture. The slurry may be dried in a vacuum oven at about 50 to 100 °C, such as 80 °C, for a period of time, say 6 h, 12 h, 18 h, 24 h, 30 h, or more.

[0044] After drying and obtaining the powder mixture, the powder mixture is heated until the powder mixture is melted. In various embodiments, the powder mixture is melted at a temperature range of about 1,000 to 1,500 °C, such as at about 1,375 °C. The heating and melting step may be carried in more than one step. For example, the powder mixture may be first heated in an alumina crucible at about 375 °C for about 2 h, followed by re-grinding the synthesized powders and melting the grounded synthesized powders at 1,375 °C for another 2 h in a platinum crucible.

[0045] The melt is then rapidly quenched. For example, the viscous melt may be poured onto stainless-steel plates preheated between about 150 °C and 500 °C, such as between 200 °C and 400 °C, between 250 °C and 350 °C, or at about 300 °C and quenched in a quenching medium immediately. In one embodiment, the powder mixture is quenched on a stainless steel plate preheated at about 300 °C.

[0046] After quenching, the quenched powder mixture is annealed at a temperature range of about 500 to 600 °C, such as about 550 °C to obtain a lithium-aluminium-germanium-phosphate (Li₂0- Al₂0₃-Ge0₂-P₂0₅) glass.

[0047] Following the annealing step, the Li₂0-Al₂0₃-Ge0₂-P₂0₅ glass is doped with B₂0₃ to obtain the Li₂0-Al₂0₃-Ge0₂-P₂0₅ electrolyte system. In various embodiments, the Li₂0-Al₂0₃- Ge0₂-P₂0₅ glass is doped with at most about 0.5 wt% B₂0₃, based on the total weight of the electrolyte system. For example, the glass ceramic electrolyte system contains about 0.5 wt% B₂0₃ (i.e. LAGP-0.5B₂O₃), about 0.4 wt% B₂0₃ (i.e. LAGP-0.4B₂O₃), about 0.3 wt% B₂0₃ (i.e. LAGP- 0.3B₂O₃), about 0.2 wt% B₂0₃ (i.e. LAGP-0.2B₂O₃), about 0.1 wt% B₂0₃ (i.e. LAGP-0.1B₂O₃), or less.

[0048] In one embodiment, the glass ceramic electrolyte system contains about 0.3 wt% B₂0₃ (i.e. LAGP-0.3B₂O₃).

[0049] For the present LAGP-xB₂0₃ (x = 0-0.4 wt.%) GCE system, the heterogeneous glass forming composition with considerably increased ionic conductivity is optimized. The solid-state

GCEs mitigate bridging of dendrites between lithium-metal anode and cathode. Thus, the compatibility of the optimized composition of LAGP-xB₂0₃ (x = 0-0.4 wt.%) series with lithium metal has been analyzed by evaluating the impedance variation of symmetrical Li/ LAGP-xB₂0₃ GCEs/Li cell of storage up to 9 days. It has been demonstrated that the solid-state B₂0₃-doped LAGP GCEs, are highly stable in water for more than 6 weeks as well as possess very high ionic conductivity. The B₂0₃-doped LAGP GCEs may therefore be considered as the potential candidates in future lithium/air batteries.

[0050] Accordingly, a third aspect of the invention relates to a lithium-ion battery, comprising a glass ceramic electrolyte system sandwiched between an anode and a cathode, wherein the glass ceramic electrolyte system comprises a lithium-aluminium-germanium-phosphate (Li₂0-Al₂0₃- Ge0₂-P₂0₅) electrolyte doped with boron oxide (B₂0₃).

[0051] In various embodiments, the Li₂0-Al₂0₃-Ge0₂-P₂0₅ glass is doped with at most about 0.5 wt% B₂0₃, based on the total weight of the electrolyte system. For example, the glass ceramic electrolyte system contains about 0.5 wt% B₂0₃ (i.e. LAGP-0.5B₂O₃), about 0.4 wt% B₂0₃ (i.e. LAGP-0.4B₂O₃), about 0.3 wt% B₂0₃ (i.e. LAGP-0.3B₂O₃), about 0.2 wt% B₂0₃ (i.e. LAGP- 0.2B₂O₃), about 0.1 wt% B₂0₃ (i.e. LAGP-0.1B₂O₃), or less.

[0052] In one embodiment, the glass ceramic electrolyte system contains about 0.3 wt% B₂0₃ (i.e. LAGP-0.3B₂O₃).

[0053] In order that the invention may be readily understood and put into practical effect, particular embodiments will now be described by way of the following non-limiting examples. Examples

[0054] Effect of B₂0₃ doping on structural, morphological and conductivity of superionic LAGP GCEs system have been investigated by various techniques such as differential scanning calorimetry (DSC), X-ray diffraction (XRD), field-emission scanning electron microscope (FE- SEM). The relationship between thermal properties and glass structure is also discussed. The local structure i.e. Li dynamics of the LAGP-xB₂0₃ (x = 0.0-0.4 wt.%) together with the Li-0 coordination in GCEs was investigated by solid-state ⁷Li NMR spectroscopic technique. In order to describe the various sites of the Li in the NaSICON-like structure, chemical shift data expressed against solid 1M LiCl are used. The ionic conductivity of the LAGP-xB₂0₃ (x = 0-0.4 wt.%) glass system was studied in the frequency range of 100 mHz to 1 MHz in the wide temperature range -30 - 100 °C. Hence, the addition of B₂0₃ incorporated in the parent LAGP GCEs was: (i) to reduce the crystallinity: (ii) to suppress A1P04 impurity phase and inhomogeneity; and (iii) to improve the ionic conduction resulting from the increased cation concentration in the glass system.

Experimental

[0055] Synthesis of glass powder, green pellet, and sintered disks. Various contents of reagent- grade Li₂C0₃, A1₂0₃, Ge0₂ and ΝΗ₄Η₂Ρ0₄ in a stoichiometric ratio for preparing

Li] .5Al₀.5Gei.5(PO₄)₃ (LAGP) were milled thoroughly in a horizontal planetary mill for 24 h in alcohol (10 vol. % solid contents). It is important to be noted that because of the volatile nature of lithium compounds at high melting temperature and in order to maintain the stoichiometry, an extra amount of lithium compound has been added to the mixture. The mixed slurries were dried in a vacuum oven at 80 °C for 24 h. After that, 20 g powder mixture was heated in an alumina crucible at 375 °C for 2 h. The synthesized powders were reground and then melted at 1,375 °C for 2 h in a platinum crucible. The viscous melt was poured onto stainless-steel plates preheated at ~ 300 °C and quenched immediately. Within the glass-forming range, these glasses were colourless and transparent and subsequently annealed at 550 °C for 2 h to release the glass stresses. Thereafter, reagent-grade B₂0₃ has been mixed to the parent LAGP glass in the different weight ratios to analyse the doping effect. Minimum dopant level is necessary to stabilize the NaSICON- rhombohedra (R-3cH) phase and to suppress the impurity A1P0₄ phase. The thus-obtained glass frits were then mixed with different amount of B₂0₃ and ground to fine powder by high energy ball milling for 3 h to obtain particle size smaller than 75 μν . Green glass "pellets" of 16 mm diameter and - 1-2 mm thickness were made by die-pressing 400 mg glass powder in a stainless-steel die at a pressure of 6 tons/cm2 held for 60 s. The pellets were sintered at 825-875 °C for 12 h with a programmed heating and cooling rate of 3 °C/min in air.

[0056] Materials Characterization. In order to identify the characteristic temperatures of the fine- powdered glass, thermal analysis was conducted by a differential scanning calorimeter (TA

Instruments Model 2910). Specimens were scanned at the rate of 10 °C/min in a temperature range of 30 to 900 °C under nitrogen atmosphere, and air-flow at 50 ml/min. The 30 mg specimen was placed in an alumina pan and an empty alumina pan was used as a reference for the measurement. Based on present DSC data, thermal treatment was carried out in two ways. In the first step, the samples were heated to the temperature between glass transition temperature (Tg) and

crystallization temperature (Tc) for 2 h in order to achieve nucleation. In the second case, the glassy samples were heated at a constant temperature, above Tc for the same duration of 12 h to obtain the GCEs.

[0057] The crystalline phases of the sintered GCEs were identified by X-ray diffraction (XRD) technique using Bruker D8 AXS Advance X-ray diffractometer with CUKQ radiation (λ = 1.5406 A) with step scanning (0.02°, 0.6 s dwell time, 40 kV) over a 2Θ range of 10 to 80°. To determine the lattice parameters as well as the quantitative phase composition, further evaluation of the diffraction patterns by means of the Rietveld method was carried out using the TOPAS software. Polished cross-section and plane- view of the samples after heat treatment were characterized by high resolution field-emission scanning electron microscope (FE-SEM: JEOL JSM7600F) at an accelerating voltage of 5 kV. The samples were mounted on metal stubs using conductive double- sided carbon tape, and a thin layer of platinum was sputtered on the sample using JEOL JFC-1200 prior to scanning. X-ray energy dispersive spectroscopy (EDS, DX-4, EDAX Co., USA) of the samples were recorded with computer interfaced INC A mapping software attached with FE-SEM to investigate crystalline features and the microstructural changes of the interface.

[0058] To gain the insights to the microscopic ion dynamics, solid-state ⁷Li static nuclear magnetic resonance (NMR) measurements were performed at 1 16.6 MHz on a Bruker DSX-300

spectrometer (Rheinstetten, Germany) operated by a Tecmag interface. A recycle pulse delay of 2 s and 64-320 scans were used. At room temperature, quadrupole interactions remain small (CQ<180 kHz), making the irradiation of the central and satellite transitions non selective. Magic Angle Spinning (MAS) was performed using a cylindrical zirconia spinner rotated at about 5 kHz. ⁷Li NMR chemical shifts were given relative to 1M LiCl solid solution. The fitting of the NMR spectra was carried out with the Bruker WINFIT software package. This program allows the position, full width at half maximum (FWHM), which is commonly referred to as the 'line width' and intensity of NMR peaks to be determined. However, quadrupole CQ and η values have to be deduced with a trial and error procedure.

[0059] For conductivity measurements, gold layers were sputtered on both surfaces of the sintered

LAGP-xB₂0₃ (0-0.4 wt.%) pellets (dia ~ 16 mm, and thickness - 1-2 mm). Ionic conductivity measurements were carried out with an AC impedance method using Solartron 1470E and SI

1255B Impedance/gain-phase analyzer coupled with a potentiostat, in a frequency range from 10 mHz to 1 MHz at an AC amplitude voltage of 100 mV in the temperature range of -30 - 100 °C. The conductivity measurements were recorded during the cooling cycle assembled in stainless steel (SS) Swagelok® cells. The cell was kept at each measuring temperature for a minimum of 30 min to ensure thermal equilibration of the sample at the temperature before measurement.

[0060] Stability tests. The time dependant interfacial resistance (Rf) of B₂0₃ doped LAGP GCEs, sandwiched between two lithium electrodes to form a symmetrical Li/GCEs/Li cell was measured at room temperature by measuring their impedance response. The cell was assembled and sealed in an argon- filled MBraun glove box (02 < 0.1 ppm; H20 < 0.1 ppm). Impedance spectra were collected for 9 days at ambient temperatures. Samples of LAGP-xB203 (0-0.4 wt.%) were soaked in deionized water in the ambient environment at room temperature for 6 weeks. The stability is determined from the weight change and physical appearance of the GCEs.

Results and discussion

[0061] Thermal characterization ofLAGP-xB Os (0-0.4 wt. %) GCEs. Differential scanning calorimetry (DSC) was performed on LAGP-xB₂0₃ (x = 0-0.4 wt.%) under nitrogen-atmosphere to investigate the crystallization behaviour and kinetic factors and, thermograms are displayed in Fig.

1. Significant differences in thermal behaviour have been observed for different values of x, which can be attributed to degree of crystallization. The glass transition temperatures (T_g) and

crystallization temperatures (T_c) are estimated for all samples. T_g is found to decrease with increasing x with a minimum T_g achieved for x = 0.3 wt.% at 497 °C, which is about 26.3 °C lower than the parent LAGP glass. For x = 0.4 wt.% composition, T_g is slightly higher than that observed for x = 0.3 wt.% composition. From the DSC data, for the parent LAGP glass, T_g and T_c of the annealed glasses were determined to be 523.3 and 701.9 °C, respectively. These values are about 24.6 and 48 °C lower than those observed by others, respectively. This may be due to different processing parameters and melting conditions. The exothermic peaks on the DSC profiles are resulted from the crystallization of the glassy matrix, the nature of as-quenched materials. T_c is observed to be maximum for x = 0.3 wt.% at 748.6 °C, which is about 47 °C higher than the parent LAGP glass, while minimum T_c is observed for x = 0.4 wt.% composition. As can be seen, T_c increases gradually with increasing x, suggesting that the crystallization of the glasses is difficult at higher values of x and indicates that the excess B₂0₃ impedes the crystallization of the phosphate glass. It is clear that the values of T_c-T_g are observed to vary with increasing x, and found to be maximum for x = 0.3 wt.% (maximum thermal stability) and minimum for x = 0.4 wt.% (minimum thermal stability) indicating that the thermal stability of the glasses against crystallization vary with B₂0₃ doping in the LAGP glass matrix. In other words, excess B₂0₃ doping has little influence on the thermal stability of the glasses in the LAGP system.

[0062] Structural/Morphological characterization of GCEs. The XRD patterns of LAGP GCEs sintered at temperatures 825-875 °C for 12 h (as shown in Fig. 2) confirm the crystalline

NASICON-type phase formation. The parent LAGP GCEs system indicates the formation of a solid solution as Al³⁺ ions replace the Ge⁴⁺ ions. The refinement of the LAGP crystal structure performed by a Rietveld analysis of the powder X-ray diffraction patterns shows a rhombohedral symmetry (space group R-3c, #167). In the LAGP phase, the lithium is located in the partially occupied Ml site (Wickoff site 6b), while Al³⁺ and Ge⁴⁺ share the 12c site. This structure is observed for all the samples sintered. The nominal composition and occupancy of the sites were kept constant for all refinements. Diffraction peaks at the 2Θ = 26, 37.5 and 39.5 ° due to the impurity phase berlinite-type A1P0₄ (space group C2221 # 20) appears at all sintering temperatures. The major diffraction peaks became narrower and narrower, with an increase in sintering temperature.

[0063] A slight variation in peak intensity indicates variation in the samples crystallinity, i.e., degree of structural order in a solid as well as growth of larger grains as clearly evident for the samples sintered at 850 °C for 12 h. This higher crystallinity is expected to develop better connection between the conducting channels in the NaSICON structure enhancing the material's ionic conductivity.

[0064] The XRD patterns of LAGP-xB₂0₃ (x = 0-0.4 wt.%) GCEs sintered at 850 °C for 12 h are shown in Fig. 3, indicating well defined peaks comparable to those observed for parent LAGP samples. The impurity phase, A1P0₄ is observed for the B₂0₃ doped samples. The peak intensity

(in contrast with the full width at half maximum) of the NaSICON-type phase decreases slightly with increasing x up to 0.3 wt. %, indicating that the excess B2O3 (> 0.3 wt.%) in the LAGP-xB₂0₃ glasses makes the crystallization process bit difficult, which is consistent with the DSC results (Fig. 1). It can also be seen that the intensity of impurity peaks decreases rapidly with increasing x up to 0.3 wt.% perhaps making the GCEs more conducting. It is clear that a part of the excess B₂0₃ (x = 0.4 wt.%) suppresses the primary LAGP as well as secondary AIPO₄ phase in the LAGP-xB₂0₃ GCEs indicating the disordered crystal structure. Refined crystallographic structural a, c

parameters and cell volume and amounts of secondary phases of sintered LAGP-xB₂0₃ GCEs, determined from the Rietveld analysis using TOPAS software, are listed in Table 1 below. The lattice parameters obtained agree well with the literature data for LiGe₂(P0₄)₃, a = 8.268 A°, and c = 20.645 A° (JCPDS 41-0034), thus confirming the NaSICON type structure of the crystalline phase for the B₂0₃ doped samples. The lattice constant a exhibits linear increase over the whole composition range. Comparison of the data reveals a slight decrease in c parameter (may be due to change in the structure/microstructure) as well as in the cell volume with increasing x from 0 to 0.3 wt.%, which is probably caused by a change of the 3D network structure.

Table 1 - Rietveld Analysis Data and Unit Cell Parameters for LAGP-xB;0₃ (x

%) GCEs sintered at 850 °C for 12 h.

[0065] Micromorphology and element distribution of fractured surface of LAGP-xB₂0₃ (0-0.5) glass-ceramic electrolytes were analyzed by field emission-scanning electron microscope (FESEM) equipped with energy dispersive spectrometry (EDS). FE-SEM micrographs of LAGP heat-treated at various temperatures 825-875 °C as shown in Fig. 4 showed a uniform surface morphology and the well-crystalline grains network with fully interconnected grain boundaries. LAGP GCEs sintered at 875 °C shows enhanced crystal growth thus higher density than in case of those sintered at 825 °C. Highly crystalline grains (1-10 /mi) are observed for parent LAGP GCEs. The GCEs heated at 850 °C for 6 h was characterized by very dense structure with very few pores and cracks. When the sintering temperature was further increased to 875 °C, the crystallized grains in the glass- ceramics grew further; moreover some cracks occurred along the grain boundaries. Fig. 5 shows the FESEM images of fracture sections for LAGP-xB₂0₃ glass-ceramics treated at 850 °C for 12 h. It showed dense structure without any large defects such as cracks or voids for all compositions.

Significant differences between the crystal morphology as a result of different B₂0₃ doping are observed owing to several factors, including defect density and degree of impurities in parent LAGP lattice. The density of the GCEs was determined by the Archimedes method with ethanol as the immersion medium. The crystal morphology of parent LAGP glass was rather dense (density ~ 95-97 %) and smooth as compared to B₂0₃-doped LAGP glass (density ~ 93-96 %). The grains are in good contact with each other although a certain amount of pores in the boundary area can be observed. The variation in ionic conductivity may not solely depend on the variation in

microstructure and preparation methods. This distinction in conductivity may arise because of several factors, including bulk and grain-boundary contributions as a result of particle to particle contact, density, sintering, and electrode-electrolyte interface effect. The distribution of Ge and Al elements obtained by EDS is shown in Table 2. The Ge and Al elements seem to vary evenly with B₂0₃ content.

Table 2 - EDX Analysis of LAGP-xB₂<¼ (x = 0.0-0.4 wt. %) GCEs sintered at 850 °C for 12 h.

[0066] ⁷ Li NMR spectroscopy studies on LAGP-XB2O3 (0-0.4 wt %). The structural sites occupied by lithium for LAGP-xB203 (0-0.4 wt.%) GCEs are determined by ⁷Li NMR (I = 3/2) spectra (shown in Fig. 6) recorded in static conditions at room temperature. In the inset of Fig. 6, the broad view of the central component of the ⁷Li static NMR spectra of the series is shown. The spectra are formed by a central line (-1/2, 1/2 transition) and two satellite lines (-3/2, -1/2 and 1/2, 3/2 transitions). Central and satellite lines are modulated by equally spaced spinning sidebands, the powder patterns of samples being reproduced by connecting the tops of the sidebands. The correlation between ⁷Li line width and ionic conductivity is investigated by analyzing the full width at half maximum (FWHM) of the central ⁷Li transition as a function of B₂0₃ doping (in different wt.%) in parent LAGP matrix. The FWHM of the central transition is near 220 Hz in the static ⁷Li

NMR spectra for the parent LAGP matrix. The line width of this component is almost constant for x <0.2 but increases for x > 0.2, which aligns with the conductivity trend (Fig. 9 and 10). At the same time, a progressive shift of this component from 1 to 3.5 ppm is observed. To the first order,

⁷Li central transition line widths are governed mainly by two effects: (i) the strength of the dipole- dipole interaction (most likely Li-P) and (ii) the lithium ion mobility (correlated with conductivity) which can effectively average these interactions. The latter effect is most likely the main cause of line broadening in the LAGP-xB₂0₃ (0-0.4 wt.%) GCEs system because conductivity is found to increase drastically for x - 0.3 wt.% as will be discussed below. The increase in the strength of the dipole-dipole interactions results from an increase or stretching in the 3D tunnel framework responsible for Li ion migration, which therefore validates the incorporation of a part of the excess lithium into the crystal lattice of the NaSICON-type structure. Because of a large overlap of ⁷Li

NMR signals, L1O4 and Li0₆ units are not well distinguished in the presently investigated LAGP- xB₂0₃ (0-0.4 wt.%) GCEs system. Spinning sideband patterns were used to determine the quadrupole constant, C_Q and the asymmetry parameter, η of the LAGP-xB₂0₃, both of which characterize the structural sites occupied by Li. Above x = 0.3 wt.%, the number of spinning sidebands increases, indicating the presence of another disordered phase. In samples with x = 0.3 and 0.4 wt.%, the ⁷Li NMR spectra are formed by two components, with quadrupole constants C_Q

(1) 40 kHz and CQ (2) 110 kHz. Finally, an increase of the C_Q (1) parameter of NASICON compounds is observed. The sharpness of these satellite transitions appears to be dependent on the substitution of B₂0₃ as well as extent of processing, consistent with some distribution of the electric field gradient (efg) experienced by individual lithium ions. Substitution does not, however, appear to have a major effect on the magnitude of the efg, as seen in Fig. 6, where only the addition of B₂0₃ appears to cause a very slight but steady increase in CQ (from 42 kHz to 43.2 kHz).

Previous investigation of ⁷Li in LGP have suggested that this value of C_Q corresponds to lithium being present predominantly in the Ml site in the NaSICON structure; that is, the site in which lithium ions are coordinated by six oxygen in octahedral arrangement and in which the efg is relatively symmetric, as compared with the M2 site. The slight increase in C_Q upon addition of

B₂0₃ may be a result of some Ml substitution as well as contraction of the overall cell volume. [0067] Ionic conductivity of LAGP-XB2O3 (0-0.4 wt. %) GCEs. The total ionic conductivity of NaSICON-type electrolyte materials depends on Li ion transport in the crystalline grains as well as through the grain boundaries. As known, those grains are more conducting than grain boundaries, thus controlling the overall conductivity of the polycrystalline material embedded in an

electrochemical cell. Therefore, by controlling the structure/microstructure and the Li ion transport properties through the grain boundaries is a key factor for achieving superionic GCEs.

[0068J The complex impedance plots (-Z" vs. Z') recorded at room temperature for LAGP sintered at 825-875 °C for 12 h are displayed in Fig. 7. One arc followed by an inclined spike at the low frequency end of the plot is observed for all samples. The disappearance of the bulk impedance semicircle in the impedance profile at room temperature could be ascribed to the very low resistance of the bulk, which usually responses at very high frequencies, for example, over 1 MHz. The high-frequency arc, which is observed in the frequency range 103- 105 Hz, is ascribed to the movement of Li ions inside the grains (grain-interior response). The spike, which is observed in the range 100 mHz - 1 kHz, is due to the blocking effect of the ions at the electrode surfaces (electrode response); this response is better observed at higher temperatures.

[0069] The total resistance (Rb + R_gb) of the sample is obtained from the right intercept of the semicircle with the real axis in the plots. The value of the bulk resistance (R_b) is therefore the difference between right intercept from the left intercept of the semicircle with the real axis. The R_b value is used to calculate the corresponding values of a (total conductivity) for all the samples. In the case of the LAGP GCEs, only one semicircle is observed corresponding to a bulk conduction mechanism since there are grain boundaries present in the crystalline LAGP GCEs. It is noteworthy that the Rb as well as Rb + R_gb are observed to be minimum for the GCEs treated at 850 °C, whereas R_gb is slightly higher as a result of the extra growth of crystalline material. R_b is observed to be the minimum ~ 49 Ω for the sample sintered 825 °C for 12 h. This phenomenon was consistent with those observed for ion conductivity and diffusion in ultrafine materials because the grain boundary dominates the electric transport property due to larger concentration and higher mobility of ions in the intergranular region.

[0070] The total ionic conductivity is further improved through optimizing the doping B₂0₃ element and content by: (a) increasing the density of the bulk material to minimize the grain boundary resistance; (b) introducing more Li into the framework to increase the mobile-ion concentration; and (c) tuning the lattice parameter to optimize the size of the interstitial channel for

Li ion conduction. Similar impedance plots are observed for the LAGP-xB₂0₃ (0.1-0.4 wt.%) as shown in Fig. 8. The shape of the impedance spectra (shown in an expanded scale in the inset of the figure) is consistent with other NaSICON-type electrolytes reported elsewhere. However, due to the dense structure but with interconnected voids of the sintered pellet as shown in Fig. 5, R_gb is higher than the R . Since it is difficult to separate the bulk and grain-boundary contributions at temperatures above room temperature with the impedance spectra, the total ionic conductivity was determined. It can be seen that ¾ + R_gb and Rb decreased when x increased from 0.1 to 0.3, and then rose with further increase of x. The value of Rb and R_b + R_gb reduced by a factor of 3 with the initial increase in x, and reached a minimum at around x = 0.3. The increase in R_b + R_gb and R with x > 0.4 may be probably due to the over-enlarged cell volume as obtained above and the appearance of pores and voids in intergranular region, which influenced Li ion conduction.

However, at high temperatures, the interpretation of impedance data becomes more complex because of the high-frequency imaginary impedance values becoming positive due to experimental limitations.

[0071] The temperature dependence of total ionic conductivities (σ) and activation energy (E_a) calculated from the Arrhenius equation σΤ = A (-E_a/k_bT) for parent LAGP GCEs sintered at temperature 825-875 °C for 12 h in the temperature range -30 to 100 °C is shown in Fig. 9, where A is the pre-exponential factor, E_a is the activation energy for conduction, k is Boltzmann's constant and T is the absolute temperature. The total pellet conductivity, σΤ, is obtained from the intercept of the spike and/or the arc (low-frequency end) on the Z' axis. Consequently, the total ionic conductivity is about 3.2 10^"3 S cm ¹ at room temperature, still several times higher than that of pure LiGe₂(P0₄)3 ~ 10^"5 S cm^-1 at 300 °C reported for related lithium superionic conductors with the NaSICON structure. [0072] The Arrhenius plots show two areas: (i) a straight line from -30 to 40 °C with higher activation energy ~ 0.59 eV, (ii) a change of slope around 40 °C with lower activation energy -0.21 eV. The activation energy gives a measure of the hindrance in the movement of Li ions along the conduction channels, then two regimes of activation energy are clearly differentiated as follows: the size of the bottleneck is less than that of Li ion in the first regime (below 40 °C) and larger in the second one. Samples prepared at 850 °C showed a slightly higher bulk and total Li ion conductivity compared to that of samples obtained at 825 and 875 °C. The increase in the ionic conductivity was explained due to an increase in the size of the particles from a high sintering temperature and better particle to particle contact. It also could be due to change in the mobile Li ion concentrations as a result of sintering at elevated temperatures.

[0073] Fig. 10 shows the temperature dependence of the ionic conductivity of the LAGP-XB2O3

(0.1-0.4 wt.%) sintered at temperature 850 °C for 12 h in the form of Arrhenius plots. The sample with B₂0₃ composition 0.3 displays the highest σ conductivity in the whole range of temperatures investigated in this study, attaining values in the order of 8.25 x 10^"3 S cm^"1 at 30 °C. The data obtained during the cooling cycle also follow two regions of activation energies. The B₂0₃ analogue LAGP-B₂0₃ (0.3 wt.%) exhibits a higher by a factor of three than the corresponding

LAGP GCEs, especially at low temperatures. The obvious increase in σ may probably due to the dense intergranular region, which caused the increase in the total Li ion conductivity of the specimen. The much lower conductivity of LAGP-0.4B₂O₃ is ascribed to changes in the network structure that lead to a decrease of the concentration of large voids in the structure with increasing x (i.e., no large voids at x = 0.3). Because the "window" for mobility diminishes between the occupied small "cage" cavities relative to the large "cage" sites, this structural change is clearly deleterious for Li diffusion. Furthermore, with increasing x, the Li-0 interactions in the channel become stronger, which also leads to a lowering of the conductivity. In most cases, two regions of activation energy (0.58 - 0.74 eV) and (0.18 - 0.42 eV) for lithium transport are observed. It is clear from the aforementioned results that the conductivity results are observed to be in well-agreement with the NMR results observed for the all samples.

[0074] Interfacial resistance of L1/LAGP-O.3B₂O₃ GCEs/Li symmetric cell. Fig. 11 shows the impedance response of LAGP-0.3B₂O₃ GCEs symmetric cell formed by sandwiching sample between two lithium metal electrodes at room temperature. The impedance measurements were taken at subsequent times of storage. The response evolved with the expected semicircle signature, whereby the low-frequency intercept with the real axis gave the value of the Li/GCEs interfacial resistance (Rf), which includes the resistance of the passivating film on the Li electrode surface.

Whereas the bulk resistance (R_b), given by the high frequency intercept with the real axis. Such pattern is generally shown by solid electrolytes that have high ionic conductivity (i.e. low R_b value).

Typical spectra with impedance values lower than 1000 Ω in the first two weeks of storage. The initial interfacial resistance is minimum ~ 67 Ω for Li/LAGP-0.3B₂O₃ GCEs/Li cell. The observed initial increase in the impedance implies that the lithium electrode is passivated with time due to the reactivity of the lithium electrode and the GCE samples. The observed R_f values ~ 392 Ω after

9 days of storage shows that the interfacial resistance is less for the B₂0₃ doped GCEs. Relatively lower interfacial resistance as well as a lower rate of increase in resistance is observed in B₂0₃ doped GCEs. This can be attributed to the better interfacial stability offered by appropriate B₂0₃ content. The formation of a stable solid GCE interface on the lithium metal to conduct Li ions freely as well as to prevent dendrite formation is essential for future lithium-ion/air batteries.

[0075] Stability ofLAGP-xB₂0₃ (0-0.4 wt.%) GCEs in DI water. Li ion conducting GCEs are reported to have no through-holes and do not exhibit moisture permeation as well as they are stable in alkaline aqueous solutions with and without LiCl. Therefore, to confirm the stability of LAGP- xB₂0₃ (x = 0.0-0.4) GCEs sintered at 850 °C for 12 h in DI water, stability tests by measuring the weight change subsequently for 6 weeks have been preformed. Samples submerged in small covered glass Petri dishes filled with 10 ml of distilled water. Stability in DI water for 6 weeks demonstrates no physical as well as structural degradation of these electrolytes. The weight gain seems to vary in the range of 0.8-1.4 % after 6 weeks for all the samples. No discolouration and no surface texture changes are observed. It is confirmed by appearance that even after 6 weeks, samples are having the same physical dimensions as day 1. Thus, LAGP-xB₂0₃ (0-0.4 wt.%) GCEs have the potential to protect a lithium metal anode from water and moisture in the ambient environment in future lithium/air batteries.

[0076] Conclusion. High ionic conductivity in the phosphate glass ceramic electrolytes is achieved by introducing defects/disorder in the parent glass structure. The study on the present NaSICON- type samples showed clearly that the Li ion content and their sites in the structure are critical for fast ionic conduction. The ionic conductivity of the glass ceramics sintered at 850 °C for 12 h is higher than those sintered at other temperatures. A comparison of the conductivity of the LAGP- xB₂0₃ (x = 0-0.4 wt.%) GCEs indicates that the highest conductivity - 8.75 x 10^"3 S cm^"1 at 30 °C for x = 0.3 is found for the samples sintered at 850 °C for 12 h. NaSICON phase with higher structural homogeneity, i.e., crystallinity has been observed for the parent LAGP as well as B₂0₃ doped LAGP GCEs. Stability in DI water for 6 weeks demonstrates no significant physical or microstructural degradation of these GCEs. Superionic Li ion conducting solid electrolytes with the high ionic conductivity and the low activation energy for conduction may be the potential candidate for safety prospective in future rechargeable lithium/air batteries.

[0077] By "comprising" it is meant including, but not limited to, whatever follows the word "comprising". Thus, use of the term "comprising" indicates that the listed elements are required or mandatory, but that other elements are optional and may or may not be present. [0078] By "consisting of is meant including, and limited to, whatever follows the phrase

"consisting of. Thus, the phrase "consisting of indicates that the listed elements are required or mandatory, and that no other elements may be present.

[0079] The inventions illustratively described herein may suitably be practiced in the absence of any element or elements, limitation or limitations, not specifically disclosed herein. Thus, for example, the terms "comprising", "including", "containing", etc. shall be read expansively and without limitation. Additionally, the terms and expressions employed herein have been used as terms of description and not of limitation, and there is no intention in the use of such terms and expressions of excluding any equivalents of the features shown and described or portions thereof, but it is recognized that various modifications are possible within the scope of the invention claimed. Thus, it should be understood that although the present invention has been specifically disclosed by preferred embodiments and optional features, modification and variation of the inventions embodied therein herein disclosed may be resorted to by those skilled in the art, and that such modifications and variations are considered to be within the scope of this invention.

[0080] By "about" in relation to a given numberical value, such as for temperature and period of time, it is meant to include numerical values within 10% of the specified value.

[0081] The invention has been described broadly and generically herein. Each of the narrower species and sub-generic groupings falling within the generic disclosure also form part of the invention. This includes the generic description of the invention with a proviso or negative limitation removing any subject matter from the genus, regardless of whether or not the excised material is specifically recited herein.

[0082] Other embodiments are within the following claims and non- limiting examples. In addition, where features or aspects of the invention are described in terms of Markush groups, those skilled in the art will recognize that the invention is also thereby described in terms of any individual member or subgroup of members of the Markush group.

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Claims

Claims

1. A glass ceramic electrolyte system, comprising a lithium-aluminium-germanium-phosphate (Li₂0-Al₂0₃-Ge0₂-P₂0₅) electrolyte doped with boron oxide (B₂0₃).

2. The glass ceramic electrolyte system of claim 1, containing at most about 0.5 wt% B₂0₃, based on the total weight of the electrolyte system.

3. The glass ceramic electrolyte system of claim 2, containing about 0.3 wt% B₂0₃.

4. A method for preparing a glass ceramic electrolyte system comprising a lithium-aluminium- germanium-phosphate (Li₂0-Al₂03-Ge0₂-P₂0₅) electrolyte doped with boron oxide (B₂0₃), the method comprising:

milling a mixture of lithium precursor, aluminium precursor, germanium precursor, and phosphate precursor in a solvent to obtain a slurry;

drying the slurry to obtain a powder mixture;

heating the powder mixture until the powder mixture is melted;

quenching the melted powder mixture;

annealing the quenched powder mixture to obtain a lithium-aluminium-germanium- phosphate (Li₂0-Al₂0₃-Ge0₂-P₂0₅) glass;

doping the Li₂0-Al₂0₃-Ge0₂-P₂0₅ glass with B₂0₃ or a boron oxide precursor to

obtain the Li₂0-Al₂0₃-Ge0₂-P₂0s electrolyte system.

5. The method of claim 4, wherein the lithium precursor is added in an amount more than the stoichiometric ratio of the lithium precursor, aluminium precursor, germanium precursor, and phosphate precursor in the milling mixture.

6. The method of claim 4 or 5, wherein the lithium precursor is selected from the group

consisting of Li₂C0_3, Li₂0, LiOH, Li(CH₃COO), Li₂C₂0₄, LiN0₃, LiH₂P0₄, Li₂HP0₄, Li₃P0₄, Li₂B₄0₇ and a combination thereof.

7. The method of any one of claims 4 to 6, wherein the aluminium precursor is selected from the group consisting of A1₂0₃, A1(N0₃)₃, Al(OH)₃, AIOOH, A1₂(C₂0₄)₃, Al(CH₃COO)₃, A1₂(B₄0₇)₃ and a combination thereof.

8. The method of any one of claims 4 to 7, wherein the germanium precursor is selected from the group consisting of Ge0₂, GeO, GeO(OH)₂, Ge(N0₃)₂, Ge(N0₃)₄, GeC₂0₄, Ge(C₂0₄)₂, Ge(CH₃COO)₂, Ge(CH₃COO)₄, GeB₄0₇ and a combination thereof.

9. The method of any one of claims 4 to 8, wherein the phosphate precursor is selected from the group consisting of NH₄H₂P0₄, P₂0₅, P₄0₆, P₄0₇, P₄0₈, P₄0₉, PO, P₂0₆, P₂0₃, (NH₄)₃P0₄, (NH₄)₂(HP0₄) and a combination thereof.

10. The method of any one of claims 4 to 9, wherein the boron oxide precursor is selected from the group consisting of B(N0₃)₃, B(CH₃COO)₃, B₂(C₂0₄)₃, (NH₄)₂(B₄0₇), NH₄(HB₄0₇), B(H₂P0₄)₃, B₂(HP0₄)₃, BP0₄, B(A10₂)₃, B(GeO(OH)₃)₃, B₂(Ge0₂(OH)₂)₃, B(Ge(OH)₇), B(Ge(OH)₃)₃, B₂(GeO(OH)₂)₃, B(Ge(OH)₅) and a combination thereof.

1 1. The method of any one of claims 4 to 10, wherein the solvent is an alcohol.

12. The method of any one of claims 4 to 11, wherein the solid content in the milling mixture is about 10 vol.%, based on the total volume of the milling mixture.

13. The method of any one of claims 4 to 12, wherein the powder mixture is melted at a

temperature range of about 1,000 to 1,500 °C.

14. The method of claim 13, wherein the powder mixture is melted at about 1,375 °C.

15. The method of any one of claims 4 to 14, wherein the melted powder mixture is quenched on a stainless steel plate preheated between about 150 °C and 500 °C.

16. The method of any one of claims 4 to 15, wherein the quenched powder mixture is annealed at a temperature range of about 500 to 600 °C.

17. The method of claim 16, wherein the quenched powder mixture is annealed at about 550 °C.

18. The method of any one of claims 4 to 17, wherein the Li₂0-Al₂03-Ge0₂-P₂0₅ glass is doped with at most about 0.5 wt% B₂0₃, based on the total weight of the electrolyte system.

19. The method of claim 18, wherein the Li₂0-Al₂0₃-Ge0₂-P₂0₅ glass is doped with about 0.3 wt% B₂0₃.

20. A lithium-ion battery, comprising a glass ceramic electrolyte system sandwiched between an anode and a cathode, wherein the glass ceramic electrolyte system comprises a lithium- aluminium-germanium-phosphate (Li₂0-Al₂0₃-Ge0₂-P₂C>5) electrolyte doped with boron oxide (B₂0₃).

21. The lithium-ion battery of claim 20, wherein the glass ceramic electrolyte system contains at most about 0.5 wt% B₂0₃, based on the total weight of the electrolyte system.

22. The lithium-ion battery of claim 21, wherein the glass ceramic electrolyte system contains about 0.3 wt% B₂0₃.