Porous li-ion conducting garnet ceramics

Abstract

There is provided a porous Li-ion conducting garnet ceramic having a porosity of from 20% to 45%. The garnet has a formula Li7-xDyLa3Zr2O12 where x and y are 0 ≤ x ≤ 3 and 0 ≤ y ≤ 1 respectively and D is a dopant selected from Ta, Nb, Al, Sn, Ge, Si, Li, Na, and K. The garnet is particularly suited for forming a composite with a polymer electrolyte. This composite can be used in the manufacture of all solid state batteries.

Description

POROUS LI-ION CONDUCTING GARNET CERAMICS

[0001] The present application claims the priority of U.S. provisional application no. 63 / 569922 filed on March 26, 2024 and incorporated herein by reference in its entirety.TECHNICAL FIELD

[0002] This disclosure relates to the field of solid Li-ion conductors (also known as solid electrolytes), methods of making same, and their application to solid-state lithium batteries. More particularly, the Li-ion conducting materials are from the garnet ceramic family.BACKGROUND OF THE ART

[0003] Solid-state lithium-ion secondary batteries with inorganic solid electrolytes are of interest in the field of energy storage because of their high safety, reliability, and energy density. Oxide ceramic materials belonging to the garnet family with a composition of LixLa3M20i2 (where x = 5 or 7, M = Ta, Nb, and / or Zr, exemplified by Al-stabilized Li7La3Zr20i2 also known as LLZO) have been strongly advocated as promising Li-ion conductors. These garnets are of particular interest for their chemical stability in contact with lithium metal, large electrochemical stability window extending above 5 V, and mechanical strength and safety. More importantly, they exhibit high lithium-ion conductivity (up to 10-3S COT1at room temperature) when crystallized in cubic phase. Typically, such ceramic electrolyte is synthesized and sintered at a very high temperature (circa 1100-1200 °C) in order (a) to favour the formation of cubic crystal phase as opposed to the less conductive tetragonal phase, and (b) obtain dense conducting “separators” for use in solid- state lithium batteries. Regrettably, because of their rigidity, these ceramic electrolytes suffer from large interfacial resistance in contact with the respective electrodes. In contrast to the inorganic solid-state electrolytes (SSE), polymer SSEs, such as polyethylene oxide (PEO), polyvinylidene difluoride (PVDF), polyacrylonitrile (PAN), polymethyl methacrylate (PMMA), etc. could enable good contact with electrodes leading to lower interfacial resistance, thanks to their flexibility and wettability. However, they tend to have low ionic conductivity (10-5— 10-7S cm'1at room temperature) necessitating operation at elevated temperature (-70-80 °C) introducing risks due to their potential for catching fire in addition to softness which makes them vulnerable to penetration by accidental formation of dendrites by the lithium metal anode. Accordingly, improvements are still desired to overcome the deficiencies of ceramic and polymer electrolytes.SUMMARY

[0004] In one aspect, there is provided a porous Li-ion conducting garnet ceramic having a porosity of from 20% to 45%, wherein the garnet has a formula Li7-xDyLa3Zr20i2 where x and y are 0 < x < 3 and 0 < y < 1 respectively and D is a dopant selected from Ta, Nb, Al, Sn, Ge, Si, Li, Na, and K. In some embodiments, the dopant is Al. In some embodiments, the garnet has the formula Lie.iAlo.sLasZ^O^. In some embodiments, the porosity is measured by the Archimedes method. In some embodiments, the porous Li-ion conducting garnet ceramic has a thickness of less than 2 mm. In some embodiments, the thickness is less than 100 pm.

[0005] In a further aspect, there is provided a composite material comprising the porous Li- ion conducting garnet ceramic as defined herein, and a polymer electrolyte doped with a Li salt that is infiltrated in the porosity of the porous Li-ion conducting garnet ceramic. In some embodiments, the polymer electrolyte coats at least one surface of the porous Li-ion conducting garnet ceramic. In some embodiments, the polymer electrolyte coats the porous Li-ion conducting garnet ceramic. In some embodiments, the polymer electrolyte is selected from polyethylene oxide, polyvinylidene difluoride, polyvinyl alcohol, polycaprolactone, polychitosan, polyvinyl pyrrolidone, polyvinyl chloride, and / or polyimide. In some embodiments, the polymer electrolyte is polyethylene oxide and / or polyvinylidene difluoride. In some embodiments, the polymer electrolyte is present in a concentration of from 1 wt. % to 20 wt. %.

[0006] In still a further aspect, there is provided an all-solid-state lithium battery comprising an anode, a cathode and the composite material as defined herein at the interface between the cathode and the anode. In some embodiments, the composite material is provided as a separator in the battery and as an electrolyte. The cathode in the all-solid-state lithium battery can also be formulated as a composite material with the porous Li-ion conducting garnet ceramic, the polymer electrolyte and a cathode material such as lithium iron phosphate (LFP) or lithium nickel manganese cobalt oxide (NMC such as NMC622).

[0007] In an additional aspect, there is provided an electrolyte-cathode composite, comprising a bilayer wherein the composite material as defined herein forms a first layer, and a cathode material forms a second layer, wherein the cathode material covers the first layer and infiltrates the porosity of the composite material of the first layer.

[0008] In yet an additional aspect, there is provided a method for producing the porous Li-ion conducting garnet ceramic as a powder as defined herein, comprising: providing an aqueous phase comprising Zr, La, and optionally a dopant; mixing in the aqueous phase an organicpolymer with pore forming properties to obtain a mixture; hydrothermally aging the mixture to obtain an intermediate comprising crystalline La(OH)s and an amorphous Zr hydroxide; calcining the intermediate with Li to obtain the porous Li-ion conducting garnet ceramic.

[0009] In some embodiments, there is further provided the production of a scaffold by providing the porous Li-ion conducting garnet ceramic as a powder, mixing the powder with the organic polymer with pore forming properties to obtain a powder mixture; depositing the powder mixture onto a sacrificial substrate; and calcining the porous scaffold to remove the filler polymers and obtain the porous Li-ion conducting garnet ceramic as said scaffold.

[0010] In still a further aspect, there is provided for producing the porous Li-ion conducting garnet ceramic as a scaffold as defined herein, comprising: mixing a nanoscale cubic LLZO with an organic polymer to obtain a mixture; depositing the mixture onto a sacrificial substrate; and calcining the mixture to remove the organic polymer and produce a porous Li-ion conducting garnet ceramic as a scaffold.

[0011] In some embodiments, the aqueous phase has a pH of from more than 7 to up to 14. In some embodiments, the calcining is a two-step calcination, a first step at a temperature of from 400 to 800 °C and a second step at a temperature of from 1000 to 1100 °C. In some embodiments, the organic polymer is a glycol vinyl polymer.

[0012] In yet a further aspect, there is provided a method of producing the composite material as defined herein, comprising providing the porous Li-ion conducting garnet ceramic as defined herein, and infiltrating the polymer electrolyte in the porosity of the porous Li-ion conducting garnet ceramic.

[0013] In still an additional aspect, there is provided a method of producing the composite material as defined herein, comprising producing the porous Li-ion conducting garnet ceramic as defined herein with the method as defined herein, and infiltrating the polymer electrolyte in the porosity of the porous Li-ion conducting garnet ceramic.

[0014] Many further features and combinations thereof concerning the present improvements will appear to those skilled in the art following a reading of the instant disclosure.DESCRIPTION OF THE DRAWINGS

[0015] FIG. 1 is a schematic illustration of a preparation of LLZO scaffold according to one embodiment, labelled p-LLZO scaffold-1.

[0016] FIG. 2 is a flowchart of the fabrication process of the p-LLZO scaffold-1.

[0017] FIG. 3A is a scanning electron microscopy (SEM) image of precursor powders after co-precipitation and hydrothermal aging (scale bar 5 pm).

[0018] FIG. 3B is a SEM image of precursor powders after co-precipitation and hydrothermal aging (scale bar 2 pm).

[0019] FIG. 4 shows the X-ray diffraction (XRD) patterns of the porous LLZO scaffold-1 before calcination, after initial calcination at 400 °C for 2 h, and after second calcination at different temperatures from 500 to 900 °C for 6 h.

[0020] FIG. 5A is a SEM image of the as-prepared porous LLZO scaffold-1 (before calcination).

[0021] FIG. 5B is a close-up of Fig. 5A.

[0022] FIG. 5C is a SEM image of porous LLZO scaffold- 1 after calcination at 400 °C for 2 h.

[0023] FIG. 5D is a close-up of Fig. 5C.

[0024] FIG. 5E is a SEM image of porous LLZO scaffold-1 after calcination at 800 °C for 6 h.

[0025] FIG. 6A is the X-ray photoelectron spectroscopy (XPS) O 1s spectra of the as- prepared porous LLZO scaffold-1 (before calcination).

[0026] FIG. 6B is the XPS C 1s spectra of the as-prepared porous LLZO scaffold- 1 (before calcination).

[0027] FIG. 6C is the XPS O 1s spectra of the as-prepared porous LLZO scaffold-1 after calcination at 400 °C for 2 h.

[0028] FIG. 6D is the XPS C 1s spectra of the as-prepared porous LLZO scaffold-1 after calcination at 400 °C for 2 h.

[0029] FIG. 6E is the XPS 0 1s spectra of the as-prepared porous LLZO scaffold- 1 after calcination at 800 °C for 6 h.

[0030] FIG. 6F is the XPS C 1s spectra of the as-prepared porous LLZO scaffold-1 after calcination at 800 °C for 6 h.

[0031] FIG. 7A is a SEM cross-sectional image of LLZO scaffold-1 with primary porous structure obtained after calcination under air at 800 °C for 6 h.

[0032] FIG. 7B is an energy diffraction spectroscopy (EDS) mapping of the LLZO scaffold-1 showing La.

[0033] FIG. 7C is an energy diffraction spectroscopy (EDS) mapping of the LLZO scaffold-1 showing Zr.

[0034] FIG. 7D is an energy diffraction spectroscopy (EDS) mapping of the LLZO scaffold-1 showing Al.

[0035] FIG. 8 shows the XRD pattern of the porous LLZO scaffold-1 after sintering under Ar at 1080 °C for 1 hr.

[0036] FIG. 9A is a SEM image showing the top surface of the sintered porous LLZO scaffold- 1 (scale bar 40 pm).

[0037] FIG. 9B is a SEM image showing the top surface of the sintered porous LLZO scaffold- 1 (scale bar 20 pm).

[0038] FIG. 9C is a SEM image showing the top surface of the sintered porous LLZO scaffold- 1 (scale bar 10 pm).

[0039] FIG. 9D is a SEM image showing the top surface of the sintered porous LLZO scaffold- 1 (scale bar 5 pm).

[0040] FIG. 9E is a SEM image showing the cross-section of the sintered porous LLZO scaffold-1 (scale bar 1 mm).

[0041] FIG. 9F is a SEM image showing the cross-section of the sintered porous LLZO scaffold-1 (scale bar 10 pm).

[0042] FIG. 9G is a SEM image showing the cross-section of the sintered porous LLZO scaffold-1 (scale bar 5 pm).

[0043] FIG. 9H is a SEM image showing the cross-section of the sintered porous LLZO scaffold-1 (scale bar 5 pm).

[0044] FIG. 10 is a schematic of the preparation of the hybrid solid electrolyte (HSE) by infiltration of polymer into p-LLZO-scaffold-1.

[0045] FIG. 11A is a SEM image of a surface of the hybrid (PEO / LiTFSI-in-p-LLZO) solid electrolyte, labelled HSE-1.

[0046] FIG. 11 B shows a SEM image of a cross section of the HSE-1 of Fig. 11 A.

[0047] FIG. 12A shows a SEM cross sectional image that was used for EDS analysis.

[0048] FIG. 12B shows a higher magnification SEM image of a section of Fig. 12A that was analyzed by EDS (scale bar 50 pm).

[0049] FIG. 12C shows the La in Fig. 12B.

[0050] FIG. 12D shows the Zr in Fig. 12B.

[0051] FIG. 12E shows the Al in Fig. 12B.

[0052] FIG. 12F shows the C in Fig. 12B.

[0053] FIG. 12G shows the O in Fig. 12B.

[0054] FIG. 12H is a carbon-layered map of Fig. 12B.

[0055] FIG. 121 is a combined map of Figs. 12C-12H.

[0056] FIG. 13A is the N 1s XPS spectra of polyethylene oxide (PEO)-lithium bis(trifluoromethanesulfonyl)imide (LiTFSI)

[0057] FIG. 13B is the C 1s XPS spectra of PEO-LiTFSI.

[0058] FIG. 13C is the La 3d spectra of p-LLZO scaffold-1 and HSE-1.

[0059] FIG. 13D is the N 1s XPS spectra of PEO-porous LLZO (HSE-1).

[0060] FIG. 13E is the C 1s XPS spectra of PEO-porous LLZO (HSE-1).

[0061] FIG. 13F is a schematic of the interaction between infiltrated PSE (PEO-LiTFSI) and p-LLZO scaffold- 1.

[0062] FIG. 14A shows XPS spectra for p-LLZO scaffold-1 , PEO-LiTFSI PSE and HSE-1.

[0063] FIG. 14B shows the F 1s XPS spectra for PEO-LiTFSI PSE.

[0064] FIG. 14C shows the F 1s XPS spectra for HSE-1.

[0065] FIG. 15 shows a thermogravimetric analysis (TGA) chart of the weight percentage change of polymer (PEO / LiTFSI) in the p-LLZO-based HSE-1.

[0066] FIG. 16A shows Arrhenius plots of the Li-ion conductivity (as determined by EIS) of the p-LLZO scaffold-1 infiltrated with PEO-LiTFSI or only PEO referred as HSE-1 and HSE-1 (no Li salt) compared to those of LLZO alone or PEO-TFSI alone.

[0067] FIG. 16B is a magnification of the upper three plots in Fig. 16A inside the highlighted rectangle.

[0068] FIG. 17A is a schematic of Li / HSE-1 / Li symmetric cell before plating-stripping cycling.

[0069] FIG. 17B is a SEM image of the Li / HSE-1 / Li symmetric cell (scale bar 1 mm) after plating-stripping cycling.

[0070] FIG. 17C is a SEM image of the Li / HSE-1 / Li symmetric cell (scale bar 0.5 mm) after plating-stripping cycling.

[0071] FIG. 18A shows the galvanostatic curve of the Li symmetric cells at 0.1 mA cm-2current density at T 0 - 5 h.

[0072] FIG. 18B shows the galvanostatic curve of the Li symmetric cells at 0.1 mA cm-2current density at T 495 - 500 h.

[0073] FIG. 18C shows the galvanostatic curve of the Li symmetric cells at 2 mA cm-2current density at T 0 - 5 h.

[0074] FIG. 18D shows the galvanostatic curve of the Li symmetric cells at 2 mA cm-2current density at T 495 - 500 h.

[0075] FIG. 19A shows EIS plots of the Li symmetric cells before cycling and after 500 cycles at 0.1 and 0.2 mA cm-2respectively.

[0076] FIG. 19B shows a graph of the impedance of hybrid SE and interfacial resistance before and after cycling (top curve hybrid resistance and bottom curve interfacial resistance) before and after 500 cycles.

[0077] FIG. 20 is a schematic illustration of the Li / HSE / LFP (PEO-LiTFSI , C, PVDF) all-solid- state lithium batteries (ASSLB) featuring HSE-1.

[0078] FIG. 21 A is a graph showing the charge-discharge voltage profiles of different cycles obtained from the Li / HSE / LFP(PEO-LiTFSI) ASSLB at 0.1 C and 25 °C.

[0079] FIG. 21 B is a graph showing the cycling stability obtained from the Li / HSE- 1 / LFP(PEO-LiTFSI) ASSLB at 0.1 C and 25 °C.

[0080] FIG. 22 is a graph showing the cycling performance of HSE-1 at high rate, 1 C and 2 C and high temperature, 70 °C.

[0081] FIG. 23 is a schematic illustration of the preparation of mesoporous LLZO powders via the modification of the garnet synthesis method by using organics with pore forming and plasticizer properties.

[0082] FIG. 24 is a flowchart of the synthesis process to obtain mesoporous cubic LLZO powders.

[0083] FIG. 25 shows XRD patterns of the porous cubic LLZO powders obtained at different PEG 20000 / PG molar ratios.

[0084] FIG. 26A is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 1:0.5 (scale bar 50 pm).

[0085] FIG. 26B is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 1 :0.5 (scale bar 20 pm).

[0086] FIG. 26C is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 1:0.5 (scale bar 10 pm).

[0087] FIG. 26D is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 1:1 (scale bar 50 pm).

[0088] FIG. 26E is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 1:1 (scale bar 20 pm).

[0089] FIG. 26F is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 1:1 (scale bar 10 pm).

[0090] FIG. 26G is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 2:0.5 (scale bar 50 pm).

[0091] FIG. 26H is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 2:0.5 (scale bar 20 pm).

[0092] FIG. 26I is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 2:0.5 (scale bar 10 pm).

[0093] FIG. 26J is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 2:1 (scale bar 50 pm).

[0094] FIG. 26K is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 2:1 (scale bar 20 pm).

[0095] FIG. 26L is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 2:1 (scale bar 10 pm).

[0096] FIG. 26M is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 3:0.5 (scale bar 50 pm).

[0097] FIG. 26N is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 3:0.5 (scale bar 20 pm).

[0098] FIG. 260 is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 3:0.5 (scale bar 10 pm).

[0099] FIG. 26P is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 3:1 (scale bar 50 pm).

[0100] FIG. 26Q is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 3: 1 (scale bar 20 pm).

[0101] FIG. 26R is a SEM image of porous c-LLZO powders molar ratio between PEG and PG of 3:1 (scale bar 10 pm).

[0102] FIG. 27A is a graph showing N2adsorption / desorption isotherm of the porous cubic LLZO powder.

[0103] FIG. 27B shows the Barret, Joyner and Halenda (BJH) desorption pore size distribution curve for Fig. 27A.

[0104] FIG. 27C is a top view SEM image of porous LLZO powder (scale bar 30 pm).

[0105] FIG. 27D is a top view SEM image of porous LLZO powder (scale bar 10 pm).

[0106] FIG. 27E is a top view SEM image of porous LLZO powder (scale bar 5 pm).

[0107] FIG. 27F is a top view SEM image of porous LLZO powder (scale bar 3 pm).

[0108] FIG. 28A shows a photograph of a dried green film.

[0109] FIG. 28B shows a photograph of a punched green film.

[0110] FIG. 29 is a schematic of the preparation method of producing p-LLZO scaffold-2.

[0111] FIG. 30 shows a XRD pattern of the as-prepared p-LLZO scaffold-2 and powder diffraction reference pattern of cubic LLZO phase.

[0112] FIG. 31 A is a SEM image of p-LLZO scaffold-2 (scale bar 500 pm).

[0113] FIG. 31 B is a SEM image of p-LLZO scaffold-2 top-surface (scale bar 400 pm).

[0114] FIG. 31 C is a SEM image of p-LLZO scaffold-2 cross-section (scale bar 100 pm).

[0115] FIG. 31 D is a close-up SEM image of p-LLZO scaffold-2 cross-section (FIG. 31C)(scale bar 20 pm).

[0116] FIG. 32A shows a higher magnification SEM image of a section of Fig. 31 C that was used for EDS analysis of the p-LLZO scaffold-2.

[0117] FIG. 32B is the EDS of Fig. 32A showing Zr.

[0118] FIG. 32C is the EDS of Fig. 32A showing La.

[0119] FIG. 32D is the EDS of Fig. 32A showing Al.

[0120] FIG. 33 shows XRD patterns of PVDF, p-LLZO scaffold-2, and HSE-2.

[0121] FIG. 34 shows the TGA curves of the PVDF-LiTFSI PSE film and HSE-2.

[0122] FIG. 35A shows the C 1s XPS spectra of PVDF-LiTFSI polymer solid electrolyte (PSE).

[0123] FIG. 35B shows the F 1s XPS spectra of PVDF-LiTFSI PSE.

[0124] FIG. 35C shows the N 1s XPS spectra of PVDF-LiTFSI PSE.

[0125] FIG. 35D shows the La 3d XPS spectra of HSE-2 (top curve) and p-LLZO (bottom curve).

[0126] FIG. 35E shows the C 1s XPS spectra of HSE-2.

[0127] FIG. 35F shows the F 1 s XPS spectra of HSE-2.

[0128] FIG. 35G shows the N 1s XPS spectra of HSE-2.

[0129] FIG. 35H is a schematic of the interaction between infiltrated PVDF-LiTFSI PSE and p-LLZO scaffold-2.

[0130] FIG. 36 shows XPS spectra of the p-LLZO scaffold-2, PVDF-LiTFSI PSE film, and HSE-2.

[0131] FIG. 37A shows a SEM cross-section of HSE-2.

[0132] FIG. 37B shows a zoomed-out SEM of HSE-2.

[0133] FIG. 37C shows a SEM top-surface of HSE-2.

[0134] FIG. 37D shows a SEM cross section of a commercial separator for comparison.

[0135] FIG. 38A shows a SEM image of the as-prepared HSE-2 top surface used in EDS analysis.

[0136] FIG. 38B shows the EDS of Fig. 38A for La.

[0137] FIG. 38C shows the EDS of Fig. 38A for Zr.

[0138] FIG. 38D shows the EDS of Fig. 38A for F.

[0139] FIG. 38E shows the EDS of Fig. 38A for Al.

[0140] FIG. 38F shows the EDS of Fig. 38A for O.

[0141] FIG. 39 shows the Arrhenius plots of the HSE-2, cubic LLZO, LLZO-in-PVDF-LiTFSI, and PVDF-LiTFSI PSE.

[0142] FIG. 40 is schematic diagram showing the Li / HSE-2 / Li symmetric cell.

[0143] FIG. 41 is a graph showing the galvanostatic curves of the Li / HSE-2 / Li symmetric cells at 0.1 and 0.5 mA cm'2.

[0144] FIG. 42 is a magnification of the galvanostatic curves of Fig. 41 from 995 to 1000 hours.

[0145] FIG. 43 shows a linear sweep voltammetry (LSV) curve of HSE-2 at room temperature with scan rate 10 mV s-1.

[0146] FIG. 44A is a SEM image of Li symmetric cell after 1000 h cycling at 0.5 mA cm-2(scale bar 500 pm).

[0147] FIG. 44B is a SEM image of Li symmetric cell after 1000 h cycling at 0.5 mA cm-2(scale bar 200 pm).

[0148] FIG. 44C is a SEM image of Li symmetric cell after 1000 h cycling at 0.5 mA cm-2(scale bar 30 pm).

[0149] FIG. 44D shows a EDS map of the cycled interface between Li metal and HSE-2.

[0150] FIG. 44E is an EDS showing the La for Fig. 44D.

[0151] FIG. 44F is an EDS showing the Zr for Fig. 44D.

[0152] FIG. 44G is an EDS showing the Al for Fig. 44D.

[0153] FIG. 44H is an EDS showing the F for Fig. 44D.

[0154] FIG. 44I is an EDS showing the O for Fig. 44D.

[0155] FIG. 45A is a XRD pattern of the as-synthesized TiCh-coated NMC622 cathode particles.

[0156] FIG. 45B is a SEM image the as-synthesized TiO2-coated NMC622 cathode particles (scale bar 100 pm).

[0157] FIG. 45C is a close-up of Fig. 47B (scale bar 50 pm).

[0158] FIG. 45D is a close-up of Fig. 47C (scale bar 10 pm).

[0159] FIG. 45E is a close-up on a single particle of Fig. 47D (scale bar 10 pm).

[0160] FIG. 45F is a transmission electron microscopy (TEM) image of the as-synthesizedTiCh-coated NMC622 cathode particles.

[0161] FIG. 46 is a schematic of a Li / HSE-2 (PVDF-LiTFSI-in-p-LLZO) / NMC622 ASSB.

[0162] FIG. 47A is a SEM of a cross-section of the HSE-2 with Li metal and TiC>2@NMC622 cathode.

[0163] FIG. 47B is an EDS mapping of the interfaces of the HSE-2 with Li metal and TiC>2@NMC622 cathode.

[0164] FIG. 47C is an EDS showing the Zr of Fig. 47B.

[0165] FIG. 47D is an EDS showing the La of Fig. 47B.

[0166] FIG. 47E is an EDS showing the F of Fig. 47B.

[0167] FIG. 47F is an EDS showing the Mn of Fig. 47B.

[0168] FIG. 47G is an EDS showing the Ni of Fig. 47B.

[0169] FIG. 47H is an EDS showing the Co of Fig. 47B.

[0170] FIG. 47I is an EDS showing the Al of Fig. 47B.

[0171] FIG. 48A is a graph showing charge-discharge voltage profiles of different cycles for a TiO2@NMC622 / HSE-2 / Li battery with HSE-2.

[0172] FIG. 48B is a graph showing cycling stability at 0.1 C and 25 °C for the battery of Fig. 48A.

[0173] FIG. 49A shows charge-discharge voltage profiles of different cycles for conventional lithium-ion battery (NMC622 cathode / Li metal anode) with liquid electrolyte.

[0174] FIG. 49B is a graph showing cycling stability at 0.1 C and 25 °C for the battery of Fig. 49A.

[0175] FIG. 50A shows charge-discharge profiles at different rates for a battery with HSE-2 (TiO2@NMC622 / HSE-2 / Li) in the range of 3 to 4.5 V at 25 °C.

[0176] FIG. 50B shows the rate performance from 0.1 C to 2 C at 25 °C for the battery of Fig. 50A.

[0177] FIG. 51A is a graph showing cycling stability or TiO2@NMC622 / HSE-2 / Li ASSB in the range 3 to 4.5 V at 0.2C and 40 °C.

[0178] FIG. 51B is a graph showing cycling stability or TiO2@NMC622 / HSE-2 / Li ASSB in the range 3 to 4.6 V at 0.2C and 40 °C.

[0179] FIG. 51C is a graph showing cycling stability or TiO2@NMC622 / HSE-2 / Li ASSB in the range 3 to 4.7 V at 0.2C and 40 °C.

[0180] FIG. 51 D is a graph showing cycling stability or TiO2@NMC622 / HSE-2 / Li ASSB in the range 3 to 4.8 V at 0.2C and 40 °C.

[0181] FIG. 52A is a graph showing the EIS curve of the as-assembled TiO2@NMC622 / HSE- 2 / Li ASSB before CV measurements from 3-4.8 V at a scan rate of 0.1 mV s-1.

[0182] FIG. 52B is a graph showing EIS curves of the as-assembled TiC>2@NMC622 / HSE- 2 / Li ASSB after CV measurements from 3-4.8 V at a scan rate of 0.1 mV s-1.

[0183] FIG. 52C is a graph showing CV profiles at different cycles of the as-assembled TiO2@NMC622 / HSE-2 / Li ASSB.

[0184] FIG. 53A is a SEM image showing the morphology of cycled interface TiC>2@NMC622 I HSE-2 with a scale bar of 20 pm.

[0185] FIG. 53B is a close-up SEM image of region 1 in FIG53A.

[0186] FIG. 53C is a close-up SEM image of region 2 in FIG. 53A.

[0187] FIG. 54A is a rate performance plot of the as-assembled TiO2@NMC622 / HSE-2 / LiASSB in the range of 3 to 4.5 V at 40 °C.

[0188] FIG. 54B is a rate performance plot of the as-assembled TiO2@NMC622 / HSE-2 / Li ASSB in the range of 3 to 4.6 V at 40 °C.

[0189] FIG. 54C is a rate performance plot of the as-assembled TiO2@NMC622 / HSE-2 / Li ASSB in the range of 3 to 4.7 V at 40 °C.

[0190] FIG. 54D is a rate performance plot of the as-assembled TiO2@NMC622 / HSE-4 / Li ASSB in the range of 3 to 4.8 V at 40 °C.

[0191] FIG. 55A is a SEM image showing a cross section of the interface between the ceramic-based CSE (HSE-2) and the cathode TiO2@NMC622 in the battery.

[0192] FIG. 55B is a schematic of the interface section examined with focused-ion beam (FIB) and SEM technique.

[0193] FIG. 55C is an annotated SEM image of the morphology of the interface between the ceramic-based CSE (HSE-2) and the cathode TiO2@NMC622 in the all-solid-state battery (ASSB).

[0194] FIG. 55D is a SEM image of the as-assembled TiO2@NMC622 I HSE-3 I Li ASSB (before cycling) cross section.

[0195] FIG. 55E is EDS quantitative spectra of different regions of SEM image of Figure 55D before cycling of ASSB.

[0196] FIG. 55F is EDS quantitative spectra of different regions of SEM image of Figure 54D after cycling of ASSB from 3 to 4.8 V for 200 cycles.

[0197] FIG. 56A is XPS spectra of F 1s of fresh interface between the HSE-2 and TiC>2@NMC622 cathode.

[0198] FIG. 56B is XPS spectra of F 1s of cycled interface between the HSE-2 and TiC>2@NMC622 cathode.

[0199] FIG. 56C is XPS spectra of N 1s of fresh interface between the HSE-2 and TiC>2@NMC622 cathode.

[0200] FIG. 56D is XPS spectra of N 1s of cycled interface between the HSE-2 and TiC>2@NMC622 cathode.

[0201] FIG. 56E is XPS spectra of Li 1s of fresh interface between the HSE-2 and TiC>2@NMC622 cathode.

[0202] FIG. 56F is XPS spectra of Li 1s of cycled interface between the HSE-2 and TiC>2@NMC622 cathode.

[0203] FIG. 56G is XPS spectra of C 1s of fresh interface between the HSE-2 and TiC>2@NMC622 cathode.

[0204] FIG. 56H is XPS spectra of C 1s of cycled interface between the HSE-2 and TiC>2@NMC622 cathode.

[0205] FIG. 57 is a schematic illustration of the method of preparation of thin porous LLZO scaffold-3.

[0206] FIG. 58 is a X-ray diffraction (XRD) pattern of thin porous LLZO scaffold-3.

[0207] FIG. 59A is a SEM image of the as-prepared thin porous LLZO scaffold-3 (scale bar400 pm).

[0208] FIG. 59B is a close-up of Fig. 59A (scale bar 50 pm).

[0209] FIG. 59C is a close-up of Fig. 59B (scale bar 30 pm).

[0210] FIG. 59D is a cross-section SEM image of the as-prepared thin porous LLZO scaffold- 3 (scale bar 100 pm).

[0211] FIG. 59E is a close-up of Fig. 59D (scale bar 30 pm).

[0212] FIG. 59F is a close-up of Fig. 59E (scale bar 10 pm).

[0213] FIG. 60 shows X-ray diffraction (XRD) patterns of PVDF-LiFSI film, porous LLZO scaffold-3, and HSE (PVDF-LiFSI in porous LLZO scaffold-3).

[0214] FIG. 61 shows TGA curves of PVDF-LiFSI film and HSE-3 (PVDF-LiFSI-in-porous LLZO scaffold-3).

[0215] FIG. 62A is a SEM image of the as-prepared HSE-3 (PVDF-LiFSI in porous LLZO scaffold-3) (scale bar 400 pm).

[0216] FIG. 62B is a close-up of Fig. 62A (scale bar 50 pm).

[0217] FIG. 62C is a close-up of Fig. 62B (scale bar 20 pm).

[0218] FIG. 62D is a cross-section SEM image of the as-prepared HSE-3 (PVDF-LiFSI in porous LLZO scaffold-3) (scale bar 100 pm).

[0219] FIG. 62E is a close-up of Fig. 62D (scale bar 30 pm).

[0220] FIG. 62F is a close-up of Fig. 62E (scale bar 10 pm).

[0221] FIG. 63A is an EDS mapping image of the as-prepared HSE-3 (PVDF-LiFSI in porous LLZO scaffold-3) for La element.

[0222] FIG. 63B is an EDS mapping image of the as-prepared HSE-3 (PVDF-LiFSI in porous LLZO scaffold) for Zr element.

[0223] FIG. 63C is an EDS mapping image of the as-prepared HSE-3 (PVDF-LiFSI in porous LLZO scaffold-3) for Al element.

[0224] FIG. 63D is an EDS mapping image of the as-prepared HSE-3 (PVDF-LiFSI in porous LLZO scaffold-3) for C element.

[0225] FIG. 63E is an EDS mapping image of the as-prepared HSE-3 (PVDF-LiFSI in porous LLZO scaffold) for F element.

[0226] FIG. 63F is an EDS mapping image of the as-prepared HSE-3 (PVDF-LiFSI in porous LLZO scaffold-3) for N element.

[0227] FIG. 64A is XPS spectra of C 1s of PVDF-LiFSI PSE.

[0228] FIG. 64B is XPS spectra of C 1s of HSE-3.

[0229] FIG. 64C is XPS spectra of F 1s of PVDF-LiFSI PSE.

[0230] FIG. 64D is XPS spectra of F 1s of HSE-3.

[0231] FIG. 64E is XPS spectra of N 1 s of PVDF-LiFSI PSE.

[0232] FIG. 64F is XPS spectra of N 1s of HSE-3.

[0233] FIG. 64G shows a comparison between HSE-3 and the porous LLZO scaffold-3 for La 3d XPS.

[0234] FIG. 65A shows an electrochemical impedance analysis (EIS) of PVDF-LiFSI PSE.

[0235] FIG. 65B shows an EIS of HSE-3 from room temperature to 80 °C.

[0236] FIG. 65C shows Arrhenius plots of PVDF-LiFSI PSE and HSE-3.

[0237] FIG. 66A is a DC polarization graph of PVDF-LiFSI PSE.

[0238] FIG. 66B is a AC impedance graph of PVDF-LiFSI PSE.

[0239] FIG. 66C is a DC polarization graph of HSE-3 (PVDF-LiFSI in porous LLZO scaffold- 3).

[0240] FIG. 66D is a AC impedance graph of HSE-3 (PVDF-LiFSI in porous LLZO scaffold- 3).

[0241] FIG. 67 shows Galvanostatic curves of Li | HSE-3 | Li symmetric cells under 0.1 , 0.2, 0.5, and 1 mA / cm2.

[0242] FIG. 68A shows Nyquist plots of the symmetric cells before and after galvanostatic cycling under 0.1, 0.5, and 1 mA / cm2.

[0243] FIG. 68B shows the equivalent circuit of the Nyquist plot of Fig. 68A at 1 mA / cm2.

[0244] FIG. 69 is a graph showing the galvanostatic cycling of the symmetric cell at step- increased current density.

[0245] FIG. 70 is a graph showing linear sweep voltammetry (LSV) curves of HSE-3 and PVDF-LiFSI PSE at room temperature with scan rate of 10 mV / s.

[0246] FIG. 71 is a schematic illustration of composite LLZO-LFP cathode material preparation.

[0247] FIG. 72 shows X-ray diffraction patterns of commercial LFP, nanoscale LLZO, composite cathode powders after ball-milling, and after sintering at 350 °C.

[0248] FIG. 73A shows cyclic voltammetry scans of the LFP cathode at scan rates ranging from 0.01 to 1 mV / s.

[0249] FIG. 73B is a graph showing the peak current as a function of square root of scan rate of the LFP cathode.

[0250] FIG. 73C shows cyclic voltammetry scans of the composite LLZO-LFP cathode at scan rates ranging from 0.01 to 1 mV / s.

[0251] FIG. 73D is a graph showing the peak current as a function of square root of scan rate of the composite LLZO-LFP cathode.

[0252] FIG. 74A shows the XPS spectra (Fe 2p) of the LFPcathode, LLZO-LFP composite cathode after ball-milling, and LLZO-LFP composite cathode after ball-milling and sintering.

[0253] FIG. 74B shows the XPS spectra (La 3d) of the LLZO, LLZO-LFP composite cathode after ball-milling, and LLZO-LFP composite cathode after ball-milling and sintering.

[0254] FIG. 74C shows the XPS spectra (Zr 3d) of the LLZO, LLZO-LFP composite cathode after ball-milling, and LLZO-LFP composite cathode after ball-milling and sintering.

[0255] FIG. 74D shows the XPS spectra (P 2p) of the LFP cathode, LLZO-LFP composite cathode after ball-milling, and LLZO-LFP composite cathode after ball-milling and sintering.

[0256] FIG. 74E shows the XPS spectra (Li 1s) of the LLZO, LFP cathode, LLZO-LFP composite cathode after ball-milling, and LLZO-LFP composite cathode after ball-milling and sintering.

[0257] FIG. 74F shows the XPS survey of LFP, LLZO-LFP, composite cathode after ballmilling, and composite cathode after sintering.

[0258] FIG. 75 is a TEM image of the composite LLZO-LFP cathode.

[0259] FIG. 76 is a schematic illustration of infiltration of composite LLZO-LFP cathode into porous LLZO scaffold-3.

[0260] FIG. 77A is a SEM image of porous LLZO scaffold-3 after infiltration of composite cathode and PVDF-LiFSI cross-section.

[0261] FIG. 77B is a close-up of Fig. 77A.

[0262] FIG. 77C is EDS mapping of Fig. 77B for La and Fe elements.

[0263] FIG. 77D shows the linear scanning spectra of La.

[0264] FIG. 77E shows the linear scanning spectra of Fe.

[0265] FIG. 78 is a schematic illustration of ASSB (Li / HSE-3 / LFP) assembly based on HSE- 3.

[0266] FIG. 79A is a graph showing charge-discharge curves of ASSB (Li / HSE-3 / LFP) at 0.1 C under different cycles.

[0267] FIG. 79B is a graph showing the long-term cycling of ASSB (Li / HSE-3 / LFP) at 0.1 C.

[0268] FIG. 80A is a graph showing charge-discharge curves of ASSB (Li / HSE-3 / LFP) under different rates.

[0269] FIG. 80B is a graph showing the rate performance of ASSB (Li / HSE-3 / LFP).DETAILED DESCRIPTION

[0270] To overcome the deficiencies of individual ceramic and polymer electrolytes, a hybrid polymer-ceramic solid electrolyte (SE) is provided herein. Hybrid electrolytes typically are made with an addition of ceramic fillers to the polymer matrix in order to reduce the polymer’s crystallinity and hence increase its conductivity, and to increase the strength of the material. However, the fillers’ tendency to aggregate limits their effectiveness for Lewis acid-base interaction, and the resulting ionic conductivity is insufficient for battery kinetics. Furthermore, the fillers do not successfully form an interconnected reinforcement to improve the mechanical properties of the composites due to the limited amount of ceramic fillers in polymer-based hybrid SEs. Although ultra-thin garnet-based hybrid SEs have been produced with relatively improved mechanical properties, they necessarily contain a high amount of polymer. This leads to a low ionic conductivity which is insufficient to fulfill commercial requirements. To obtain a high ionic conductivity, low crystallinity, then additional polymer chains are often required, resulting in mechanically softer polymers. The present disclosure provides a porous Li-ion conducting ceramic garnet material and structure to address this problem. This material can be used to form a porous ceramic-based hybrid SE (HSE) which overcomes the drawbacks of traditional hybrid polymer / ceramic solid electrolyte configurations. Indeed, the present SE has a robust ceramic skeleton / scaffold which provides both high ionic transport capability and high mechanical strength with the addition of infiltrated conductive polymer to address its low-resistance interfacial contact with the composite cathode and lithium metal anode electrodes. It should also be noted that Li metal is electrochemically unstable in contact with organic liquid electrolytes, and it suffers from the formation of Li dendrites. This is a further reason to avoid using a liquid electrolyte in the Li battery.

[0271] To achieve the goal of all-solid-state lithium batteries (ASSLBs) with high energy density and good safety, robust hybrid polymer-in-ceramic solid electrolytes (SEs) are desired. ASSBs overcome many problems associated with liquid electrolyte batteries which have a higher risk of fire hazard, are heavier, have less energy density per volume, and require graphite as anode. ASSBs replace graphite with lithium metal for a solid-state battery without an organic solvent, which is the main contributor to the flammability risk. Solid electrolytes that replace the liquid electrolyte are for example, oxide based, sulphide based or polymer conducting based. The oxide-based materials are traditionally preferred because they are more stable. Althoughsulphides have a higher conductivity, they also have an increased cost and require a glove box environment or handling and fabrication. The issue with polymers is that only a small amount can be utilized before the interface allowing the passage of Li ion is compromised and sufficient conductivity is lost.

[0272] The present disclosure describes the design and fabrication of porous garnet ceramics as Li-ion conductive ceramic membranes, characterized by high intrinsic lithium-ion conductivity (such as more than 10-4S / cm at room temperature (RT)) and mechanical strength (such as a Young’s modulus of more than 75 GPa). The porous garnet ceramic is made of cubic phase LLZO material synthesized by organic assisted hydrothermal processing, deposition and calcination at a temperature below 1100 °C. The porous structure is advantageously infiltrated with molecularly bonded conductive polymer filling enabling improved interfacial wettability and reduced impedance with the electrodes in ASSBs for a long cycle life function.

[0273] In one aspect, there are provided porous ceramic scaffolds (also called membranes) made of cubic garnet (otherwise abbreviated as LLZO) Li-ion conductors of variable thickness for use in solid-state energy storage devices such as lithium metal batteries. LLZO is an oxide-based garnet-type Li-ion conductor. The abbreviation LLZO as used herein refers to the lithium lanthanum zirconium oxide family with general chemical formula LixLayZrzOw(for example Li?La3Zr20i2), as well as variants of this formula (for example ±0.3 in stoichiometry for each element), and / or the addition of at least one dopant element (e.g. Al). When a dopant “D” is provided, the formula can become Li7-xDyLa3Zr20i2 (where x and y are 0 < x < 3 and 0 < y < 1 respectively covering the case where the dopant is present or not when y=0). D is selected from Ta, Nb, Al, Sn, Ge, Si, Li, Na, and K. The stoichiometric ratio, is therefore Li:La:Zr = 7:3:2 or (7- x):3:2. In one embodiment, the Al doped LLZO has the formula Li6.iAlo.3La3Zr20i2.

[0274] The ceramic Li-ion conducting porous garnet (LLZO) ceramic material is characterized by a porosity of from 20 to 45 %, from 25 to 45 %, from 30 to 45 %, from 30 to 40 %, from 30 to 35 %, or from 35 to 45 %. The porous ceramic Li-ion conducting porous garnet can be in the form of porous particles, pellets, a scaffold or a membrane. The percentage of porosity as used herein may be defined as being the total volume that is not the LLZO ceramic material (i.e. “empty” space). The porosity can for example be measured by the Archimedes method.

[0275] The porous Li-ion conducting garnet ceramic may be built at variable thicknesses (micrometer to millimeter scale) by modifying the fabrication process. In some embodiments, thethickness is of less than 10 mm, less than 5 mm, less than 1 mm, less than 500 pm, or less than 200 pm, less than 100 pm, from 50 pm to 1 mm, from 50 pm to 500 pm, from 50 pm to 200 pm, from 10 pm to 1 mm, from 10 pm to 500 pm, from 10 pm to 200 pm, from 20 to 60 pm, from 30 to 55 pm or from 35 to 50 pm. As presented in the example section below, different ceramics were synthesized with different thicknesses, ranging from mm scale thickness to 100 pm or less, such as 40 pm. In all process variations, the synthesized LLZO was a cubic phase (c-LLZO) with the composition Lie.iAlo.aLasZ^O^. In the fabrication process, a polymer filler is used to promote the formation of porosity. It should be highlighted that the polymer filler and the conducting polymer fulfill two different roles. The filler polymer is used during the fabrication and is removed via the calcination. Accordingly, the polymer filler may be conducting or not conducting, with Li-ion conducting fillers being an option in the case where one is worried that there would be residue of filler polymer that remains. In the context of applying the porous Li-ion conducting garnet ceramic in ASSBs, a more compact thickness is preferred to reduce the size of the battery, for example a thickness of less than 200 pm, preferably less than 150 pm, or even more preferably 100 pm or less.

[0276] The hybrid porous Li-ion conducting garnet ceramic can be advantageously used in ASSBs by filling the porous ceramic scaffold with the dissolved polymer electrolyte in a solvent. The solvent of the polymer is allowed to dry and leave behind the solid polymer electrolyte component. The polymer electrolyte can also be used to form a coating around the porous Li-ion conducting garnet ceramic. This results in a hybrid material made of 1 to 20 wt. % of the polymer electrolyte and the rest by the porous ceramic garnet scaffold having improved physical properties (mechanical properties at small thickness) while maintaining sufficient ionic conductivity appropriate for ASSBs. The porous Li-ion conducting garnet ceramic can also act as a solid electrolyte either alone or in a composite preferably upon pore filling by conducting or nonconducting polymer infiltration.

[0277] The porous conducting LLZO scaffolds can be advantageously formulated into hybrid (polymer-in-ceramic) electrolytes by infiltration into the porous space of polymeric / liquid or gel electrolytes enabling impedance-free interfacing with the respective Li metal anode and cathode of the energy storage device. In yet another embodiment ultra-thin, bilayer porous ceramic scaffolds are provided made of cubic garnet phase (LLZO) and an overlayer of cathode composite.

[0278] The polymer electrolyte can be any polymer that can be impregnated with a Li salt to become conductive and that is soluble in a solvent in order to be able to infiltrate the polymer in the porosity. The polymer electrolyte is, for example, selected from polyethylene oxide, polyvinyldene difluoride, polyvinyl alcohol, polycaprolactone, polychitosan, polyvinyl pyrrolidone, polyvinyl chloride, and / or polyimide. Preferably, the polymer electrolyte is polyethylene oxide (PEO), and / or polyvinylidene difluoride (PVDF). The doping of the polymer electrolyte can be performed with a Li salt such as lithium bis(trifluoromethanesulfonyl)imide (LiTFSI) or other salts (such as LiTFSI etc.). In the case of LiTFSI, the LiTFSI is dissolved in organic solvents such as acetonitrile, or N-Methyl-2-pyrrolidone (NMP) and impregnated in the polymer electrolyte. The infiltrated amount of polymer electrolyte in the resulting composite varies from 1 wt. % to 20 wt. %, from 1 wt. % to 10 wt. %, or from 3 wt. % to 10 wt. % with respect to the total weight of the composite material.

[0279] One exemplary method for producing porous Li-ion conducting garnet ceramic (or porous LLZO for short) is provided herein. First, an aqueous phase (A) of zirconium, lanthanum, and aluminum is provided, the aqueous phase having a pH between 7 and 14 which can be adjusted with LiOH. To the aqueous phase, organics are added having pore forming and plasticizer properties such as a glycol vinyl polymer (e.g. polyethylene glycol and / or propylene glycol). The nature of the organics is not important as these do not form part of the end product LLZO and only serve to help promote and form the porosity in the LLZO. The obtained mixture (B) is hydrothermally aged to yield an intermediate comprising crystalline La(OH)s and amorphous Zr / AI hydroxide from the Zr / AI and the La in the aqueous phase. This is followed by recovering, washing, and milling the intermediate which is (C) pressed and subjected to (D) two-step calcination with a Li precursor at a temperature from 400 to 800 °C to obtain the LLZO cubic phase; and further (E) calcining under inert atmosphere at 1000-1100 °C to obtain a mechanically strong porous ceramic scaffold. An alternative method could be to only follow steps A, B and D.

[0280] To obtain the composites described above having a polymer electrolyte in the porosity of the LLZO and optionally coating the LLZO, the polymer electrolyte is infiltrated by an appropriate method into the porosity. Accordingly, the polymer electrolyte needs to be soluble in a solvent (regardless of the nature of the solvent whether it is an aqueous, alcoholic, or non-polar organic solvent for example).

[0281] In one embodiment, a porous LLZO scaffold is made by using the mesoporous garnet powder obtained as described above (steps A-E or A+B+D) and by (F) formulating the garnet intoa gel using a an organic polymer acting as a binder (e.g. polyvinyl butyral and a plasticizer (e.g. benzyl butyl phthalate)), and subjecting to (G) deposition by a method such as “doctor blading” or tape casting followed by (H) optionally pressing and solvent removal, and finally (I) calcination preferably under inert atmosphere at 1000 - 1100 °C. Optionally, to obtain an ultra-thin, bilayer porous LLZO scaffolds a further step (J) is performed by depositing a standard cathode slurry (e.g. CAM / Carbon / binder; CAM stands for Cathode Active Material) on the porous LLZO framework to form an overlayer (composite cathode layer), which after (K) drying and (L) polymer infiltration is pressed (M) into a composite membrane. Accordingly, the composite can be coengineered with a top layer of cathode forming a bilayer solid electrolyte / cathode structure upon polymer infiltration and pressing.

[0282] In one embodiment, a porous LLZO scaffold is produced starting from a nanoscale cubic-LLZO (c-LLZO) which can be obtained by methods described in WO2023023856. Nanoscale refers to a size of 1 pm - 50 nm as described in WO2023023856 which is incorporated herein by reference in its entirety. The nanoscale c-LLZO can be formulated into a gel with PVB, BBP and acetone as describe above, followed by subjecting it to deposition by a method such as “doctor blading” or tape casting followed by optionally pressing and solvent removal, and finally calcination preferably under inert atmosphere at 1000 - 1100 °C. In this embodiment, a fast calcination (e.g 1 h ±5%) is preferred.

[0283] In some embodiments, as explained above, the composite membrane can be produced with a cathode material into a cathode composite material. The cathode is for example a NMC622 cathode or a lithium iron phosphate (LFP) cathode. Accordingly, there is provided an ASSB with a NMC622 cathode (LiNi0.6Co0.2Mn0.2O2) and an anode made of the LLZO scaffolds described herein. The NMC622 cathode may be coated with titanium oxide. The porosity can be filled with one of the polymer electrolytes described above such as PVDF or PVDF-LiTFSI. This ASSLB has a homogeneous interfacial layer between the cathode and the anode that provides integrity and unimpeded interfacial Li-ion conduction. When manufactured, before any battery cycling has been made, the interface between NMC622 and LLZO contains negligible amounts of Ni, Co, Mn, La and Zr. As the battery is being used and it goes through multiple cycles an intradiffused zone is created containing Ni, Co, Mn, La and Zr. There is also provided a LFP composite cathode material with a LFP core and a nanoscale LLZO shell. The composite material can be obtained by mechanically mixing the component (for example by ball milling).

[0284] One advantage of the present batteries is that the anode and cathode (along with the polymer electrolyte) can be produced together as a composite material in order to further reduce the thickness of the battery. A thin (e.g. less than 100 pm) LLZO membrane as described above can be combined with the cathode and sintered together to provide a reduced overall battery thickness, along with improved mechanical integrity. The overall thickness achieved can be less than 200 pm, less than 150 pm, or less than 100 pm.EXAMPLESFabrication of Porous LLZO Scaffold-Process variant 1

[0285] The p-LLZO scaffold-1 was prepared as described in the schematic of Fig. 1 and the flowchart given in Fig. 2. Essentially first a precursor powder of LLZO is prepared via modification of a previously developed garnet synthesis method (WO2023023856). The precursor powder preparation involved the co-precipitation of La-Zr-AI hydroxide in the presence of pore-forming organic additives (e.g. polyethylene glycol (PEG) and propylene glycol (PG)) followed by hydrothermal aging. The precursor powder is afterwards pressed into pellet form and calcined to induce the formation of pores via the gasification removal of organic molecules before subsequent calcination at higher temperature which yields a mechanically robust porous scaffold of c-LLZO. In some cases, organic additives such as polyethylene glycol (PEG) and propylene glycol (PG) were added into the synthesis solution to act as porous forming agent and plasticizer respectively for the hydrothermally aged precursor material. Optionally, the precursor material comprises La(OH)s nanofibre crystallites in a matrix of amorphous La-Zr-M (where M=AI or other doping element) hydroxide and the organic components (e.g. PEG and PG). Alternatively, the above precursor powder was mixed with a Li carrying chemical like LiOH and milled together before it was pressed into flat pellets of 1 to 2 mm thick.

[0286] The pellets were subjected to low temperature (at 400 °C) calcination in air to promote volatilization of the organics and creation of porosity followed by a second calcination step at higher temperature (at 800 °C) in air again to ensure complete removal of any residual carbon from decomposed organics and promotion of sintering into a porous scaffold of cubic LLZO. To enhance the mechanical strength of the above formed scaffold, the scaffold is further calcined at 1000-1100 °C in an inert atmosphere (e.g. Ar).Fabrication of p-LLZO scaffold-1

[0287] Lanthanum (III) nitrate hexahydrate (La(NO3)3‘6H2O), aluminium nitrate nonahydrate (AI(NO3)3‘9H2O), and lithium hydroxide monohydrate (LiOH-H2O) were dissolved in deionized (DI) water at room temperature. The amount of La must be added in excess of the stoichiometry to compensate for the solubility limitation in the elevated pH range. In this system represented as suspension A, the stoichiometric molar ratio of Li: Al: La was 6.1 : 0.3: 3.0154. Due to the spontaneous occurrence of the hydrolysis reaction, milky white precipitates were formed. Meanwhile, stoichiometric amounts of zirconyl nitrate solution (ZrO(NOs)2) (35 wt. % in dilute nitricacid) was dissolved in Dl-water at room temperature. Polyethylene glycol (PEG, molecular weight (MW) of 20000 g / mol) and propylene glycol (PG) with a molar ratio equal 2 : 0.5, were added as porous forming agent and plasticizer respectively, and labelled as solution B. After the hydrolysis and dissolution reactions were completed in both systems, solution B was added dropwise into suspension A and mixed thoroughly (at this point, the pH value of this mixed system is around 6). To further complete the co-precipitation, saturated LiOH solution (3M) was added dropwise at a rate of 1 mL / min to raise the pH of the mixed solution to >pH=7 and below the saturation concentration of LiOH (i.e. [LiOH] = ca. 5.3M, pH=11). The optimal pH value is 10.5. After this coprecipitation step, the suspension was directly transferred into a 100 mL autoclave, which was then heated to 200 °C for 4 hours (h) including 2 h for ramping and 2 h for holding at T before cooled down to RT naturally. After the hydrothermal aging step, the wet precipitates were recovered via centrifugation (thoroughly washed with Dl-water and isopropanol). Then the as- prepared precipitates were dried at 80 °C in the vacuum oven overnight. SEM images of the precursor powder collected after hydrothermal aging are shown in Figs. 3A-3B.

[0288] The dried precipitates were then mixed with the designated quantity of LiOH-H2O in isopropanol by ball milling (planetary mill) using a ZrO2 container with 1 mm ZrO2 balls at 650 rpm for 10 cycles that involved 3 min grinding / 7 min resting. The ground suspension was separated from the grinding media via a rotary evaporator at 60 °C under vacuum and dried at 80 °C in the vacuum oven overnight. The completely dry precursors were cold-pressed under 10-ton for 5 min into pellets with a diameter of 13 mm and a thickness of 1.5 mm. The pellets were sintered at 400 °C for 2 h in air atmosphere to burn out the organics, then were sintered at 800 °C for 6 h (in air) to obtain the c-LLZO phase with primary porous structure. The crystalline, morphological, and chemical / elemental characterizations of the synthesized porous LLZO are presented in Figs. 4, 5A-5F, 6A-6F, and 7A-7D.

[0289] Finally, the porous LLZO pellets were sintered at 1080 °C for 1 h under an Ar atmosphere to enhance the mechanical strength as well as the porous channel uniformity. Ambient air must be avoided because a large amount of lithium is lost resulting in the decomposition of cubic LLZO and the densification of the pellet. The XRD pattern of the p-LLZO Scaffold-1 after sintering is shown in Fig. 8. No impurities were present while its pure cubic garnet phase was preserved. Microscopic images of the top surface and cross-section of the porous c- LLZO scaffold are shown in Figs. 9A-9H. The results show that the inner pores were about 1 ~ 3 microns, and the garnet grains were tightly bonded to form a 3D porous conductive network. Theporosity of the final sintered porous LLZO scaffold-1 determined through the Archimedes method was found to be -30%.Preparation of Hybrid Solid Electrolytes using p-LLZO Scaffold-1 : hybrid solid electrolyte (HSE-1)

[0290] In this experiment, the construction of hybrid polymer-in-ceramic solid electrolytes is described. The robust porous garnet (LLZO) scaffolds described above can be advantageously combined with polymer infiltration to create hybrid solid-state electrolytes offering excellent ionic conductivity and interfacial integration with the electrodes in solid-state battery devices. Such hybrid solid-state electrolyte is described here using the porous-LLZO scaffold-1 and conductive polymer polyethylene oxide (PEO)-lithium bis(trifluoromethanesulfonyl)imide (LiTFSI) as an example. Other types of polymer-Li salt blends may also be used.

[0291] To facilitate the polymer infiltration into the pores of the ceramic scaffold the polymer is either introduced in hot molten state or preferentially dissolved in an appropriate solvent (e.g. acetonitrile) with good wetting properties towards the sintered garnet ceramic scaffold. Upon cooling or solvent evaporation-depending on method of infiltration, the polymer solidifies filling the pores and forming a cap on the exterior surface of the scaffold for facile interfacing to cathode and anode.

[0292] Alternatively, the porous LLZO scaffold may be used in conjunction with liquid or gel electrolytes for other hybrid electrolyte combinations taking advantage of the robust porous ceramic structure and the wetting ease and high conductivity of liquid electrolytes.

[0293] In this non-limitative example, the hybrid solid electrolyte (HSE-1) is prepared as described in the schematic of Fig. 10. The as-prepared sintered porous-LLZO scaffold is immersed in a polymer solution of PEO in which a Li-ion salt is added and let to infiltrate into the porous space of the scaffold before the solvent is evaporated to obtain HSE-1.Fabrication of hybrid solid electrolyte HSE-1 (p-LLZO scaffold-1 / PEO-LiTFSI)

[0294] Poly(ethylene oxide) (PEO) (MW = 600,000 g / mol) and lithium bis(trifluoromethane sulfonyl)imide (LiTFSI) were dissolved in anhydrous acetonitrile with a mole ratio of EO / Li = 10. Then, the mixture is stirred for 24 h at 55 °C in a sealed bottle to obtain a homogeneous gel. In a control test, the p-LLZO scaffold-1 was infiltrated with PEO gel only without using LiTFSI in order to evaluate the ionic conductivity of the garnet scaffold itself. To prepare the hybrid SEs, the as-prepared p-LLZO scaffold-1 was immersed into PEO gel (with and w / o LiTFSI) for 48 h and dried (via solvent evaporation) in the glove box.

[0295] The exterior surface of the PEO(LiTFSI)-in-p-LLZO hybrid solid electrolyte (HSE-1) was coated with an 8 pm-thick smooth conductive polymer layer, which facilitates the low impedance interfacial contact with the electrodes in ASSLBs as illustrated in Figs. 11A-11 B. The cross-sectional microscopic image of the as-designed ceramic-based hybrid SE in Fig. 12A shows the thickness of scaffold to be around 1.4 mm, which provides strong mechanical framework for blockage of undesirable Li-metal dendrite growth. The cross-sectional EDS results of the hybrid SE (Figs. 12B-12I) and the selected area of the dash box in Fig. 12A indicate good homogeneous distribution of polymer inside and on the surface of p-LLZO scaffold-1. Moreover, both ceramic and organic phases are continuous, which are interconnected, and interspersed to form a hybrid solid electrolyte. The continuous interface between the ceramic skeleton and PSE filler is favorable for the fast transport of Li-ion, while the as-prepared p-LLZO scaffold-1 with a high Young’s modulus can resist dendrite growth and the softer PSE can improve interfacial contact. Notably, the carbon-layered map and combined elemental map of the HSE-1 cross section (Figs. 12C-12I) clearly demonstrate that the infiltrated polymer is continuous, homogeneous, and in limited abundance as opposed other hybrid polymer-ceramic electrolytes that feature large excess of the polymer component. It further shows that the as-designed HSE-1 is ceramic-based, which provides Li-dendrite suppression and high ionic conductivity at ambient temperature. This is demonstrated in the characterizations presented below.

[0296] To clarify the interaction between infiltrated polymer (PEO) and porous c-LLZO scaffold, XPS analysis was performed on the cross-section of the hybrid SE (PEO / LiTFSI in porous-LLZO scaffold-1) plus c-LLZO powder and polymer electrolyte (PEO-LiTFSI) for comparison.

[0297] In Fig. 13A, the N 1s spectrum of the PEO-LiTFSI polymer solid electrolyte (PSE) is seen to have only one signal peak, which is located at 398.88 eV (TFSI j. For the HSE-1 (PEO- LiTFSI in p-LLZO scaffold-1) (Fig. 13D), the N 1s spectrum shows a new peak appearing at 400.38 eV corresponding to La-N bonding, which is attributed to the interaction between the La atoms of p-LLZO scaffold-1 and the TFSI' group of PSE according to the principle of hard and soft acids and bases. As for the C 1s spectrum of PSE and hybrid SE shown in Figs. 13B and 13E, the peaks are located at 284.18, 286.18, 287.58, and 292.38 eV which are attributed to (- CH2-), (-OCH2-CH2-), (-COOR), and (-TFSIj. However in HSE-1 , the four peaks shift to higherbinding energy by 0.4, 0.2, 0.7, and 0.2 eV respectively, revealing again the interaction of some groups in PSE with some metal atoms in p-LLZO scaffold-1. Meanwhile, Fig. 13C shows the La 3d spectrum of the HSE-1 which displays two sets of peaks located at 854.68, 850.28 eV (La 3d3 / 2), and 837.98, 833.48 eV (La 3d5 / 2), which shift to low binding energy by 0.8, 1.4, 0.7, and 1.4 eV compared with the peaks of p-LLZO scaffold-1 , respectively, demonstrating again the bonding of the La atoms with the polymer. Thus, it was hypothesized that the robust La-N bonding between the infiltrated PSE and p-LLZO scaffold-1 results in strong and homogeneous contact of PSE inside / outside the porous structure of LLZO scaffold as presented in Fig. 13F. The XPS surveys of the p-LLZO scaffold-1 , PEO / LiTFSI (PSE), and HSE-1 are shown in Fig. 14A, and the F 1s spectrum of the PSE and HSE-1 are presented in Figs. 14B-14C. Thermogravimetric analysis (TGA) results are shown in Fig. 15 which shows that the weight percentage of polymer (PEO / LiTFSI) in the p-LLZO-based HSE-1 was only about 3%. This shows that the as-prepared HSE-1 is mostly made of the sintered body of the porous LLZO ceramic conductor guaranteeing excellent thermal stability and safety when employed in all-solid-state lithium batteries.

[0298] To evaluate the ionic conductivity of the obtained HSE-1 , electrochemical impedance analysis (EIS) was employed. The ionic conductivity results at various temperatures of four types of electrolytes: HSE-1 (p-LLZO infiltrated with PEO-LiTFSI), HSE-1 (p-LLZO infiltrated with PEO but not LiTFSI), single LLZO SE, and single PEO-LiTFSI PSE are summarized in Figs. 16A-16B. Among them, the HSE-1 with PEO-LiTFSI had the highest ionic conductivity, 0.547 mS cm-1at 25 °C. The HSE-1 with PEO only exhibited the same ionic conductivity with the pure cubic LLZO SE, i.e. 0.43 mS cm-1at 25 °C, which demonstrated the high intrinsic ionic conductivity of the p-LLZO scaffold-1. Moreover, HSE-1 (p-LLZO infiltrated with PEO-LiTFSI) had the lowest activation energy, 0.187 eV, reflecting its enhanced Li-ion migration due to molecular bonding of PEO / LiTFSI on the highly conducting robust cubic LLZO ceramic skeleton. The superior ion conducting structure of HSE-1 was further exemplified by the high Li-ion transference number (tLi+) that is 0.71 , when the respective tLi+ is only 0.15 for PEO-LiTFSI PSE and ~ 0.25 in state of the art LLZO-in-PEO(LiTFSI) hybrid configurations. High tLi+ is a highly desirable property in solid electrolytes as it enables lower concentration polarization and lower driving force for unwanted Li dendrite growth in operating ASSLBs.

[0299] The interfacial stability and dendrite inhibition ability of the HSE-1 was evaluated using symmetric cells. The schematic image and SEM images of the Li / HSE / Li symmetric cell are shown in Figs. 17A-17C. Long-term interfacial stability is important for the operation of batteries, and this can be analyzed by constant current cycling tests as shown in Figs. 18A-18D. The results indicatethat the HSE-1 did not short-circuit at current densities of 0.1 and 0.2 mA cm-2over a duration of 500 h. Even at a higher current density of 0.4 mA cm-2, the symmetric battery still worked properly showing no increase in polarization. The SEM images in Figs. 17B-17C of the cycled Li symmetric batteries at 0.1 and 0.2 mA cm-2over 500 h provide evidence of the good interfacial contact between HSE-1 and Li metal. To further demonstrate the formed stable interface, the EIS plots of the Li symmetric cells before and after long-term plating-stripping cycling are shown in Fig. 19A and the corresponding impedance and interfacial resistance are presented in Fig. 19B. After the long-term plating-stripping cycling, the interfacial impedance only slightly increased. Therefore, the results demonstrate that the Li symmetric batteries assembled with the HSE-1 have very good cycling stability. Combine this behaviour with the mechanical strength of the ceramic scaffold it makes the present hybrid electrolyte, HSE-1 , a desirable and advantageous material to build ASSLBs with increased energy density and safety.

[0300] To demonstrate the feasibility of HSE-1 in practical applications, an ASSLB was assembled and electrochemically tested. The schematic diagram of the ASSLB assembled with the HSE-1 is presented in Fig. 20. The cathode electrode is LiFePCL / PEO-LiTFSI / Carbon / PVDF and the anode electrode is lithium metal. The results of the long-term cycling of ASSLB under 0.1C at 25 °C are shown in Figs. 21 A and 21 B, where it can be observed that the capacity is up to 163 mAh g-1and remains stable for 50 cycles, with a capacity retention rate of 98% and an average Coulombic efficiency of more than 99%. Hence, the as-designed ASSLB delivers a favorable electrochemical performance with high capacity owing to the high conductivity of HSE- 1 and the stable interfaces with both electrodes endowed by the molecularly bonded PEO-LiTFSI on the ceramic scaffold.

[0301] To further explore the stability of the as-designed ASSLB, cycling was also done at higher rates and elevated temperature, 70 °C. The results are shown in Fig. 22. As it can be seen in Fig. 22, the ASSLB exhibited good capacities at the high current rates 1 C and 2 C of 138 mAh g-1and 130 mAh g-1respectively. This suggests that the newly designed HSE-1 hybrid electrolyte enjoys high cycling and thermal interfacial stability - very important properties for use in energy storage applications.Fabrication of Porous LLZO Scaffold-Process variant No. 2

[0302] Another variation of the above fabrication method that favours engineering of thin (pm scale) porous cubic LLZO scaffold (p-LLZO scaffold-2) was developed. The scaffold wasfabricated starting from mesoporous cubic-LLZO powders. The synthesis of the latter was accomplished through the use of organic additives (with pore forming and plastizer properties) in combination with the method as described above for variant 1. The said mesoporous c-LLZO powders were then formed into high-strength porous ceramic scaffolds (p-LLZO-scaffold-2) via formulation into gel, doctor blading, pressing, and high-temperature calcination (ca. 1100 °C). The present method produced thin ceramic membrane (10-500 pm) in contrast to the first method (p- LLZO scaffold-1) that produced thick ceramic membranes (0.5-5 mm).

[0303] The p-LLZO scaffold-2 was prepared using mesoporous cubic LLZO powders. For the synthesis of the latter organic agents with pore forming and plastizer properties are used in conjuction with the garnet synthesis method as described above. A simple schematic is shown in Fig. 23, and a preferred synthesis flowchart is given in Fig. 24. The flowchart of Fig. 24 differs from that of Fig. 2 as in this case porous c-LLZO powder is produced while in Fig. 2 c-LLZO was obtained as porous scaffold following pressing+calcination of the hydrothermal precursor. In both cases polyethylene glycol (PEG) was used as pore-forming organic and propylene glycol (PG) as plasticizer.

[0304] Organic additives such as polyethylene glycol (PEG) and propylene glycol (PG) may be added in different ratios or with different molecular weight. The addition may be made in Solution A or Solution B (see Fig. 24) but the latter is preferable. PEG of different MW may be used alone or in mixture such as 2000, 6000, or 20000 g / mol. The molar ratio between PEG and PG may vary from 1 / 0, to 1 / 0.5, such as 1 / 1 , 2 / 0.5, 2 / 1 , 2 / 2, or 3 / 1. Preferably, PEG (MW 20000 g / mol) and PG are added in 2 / 0.5 molar ratio. In some cases, the c-LLZO powders are formed into high-strength porous ceramic scaffolds (p-LLZO-scaffold-2) via formulation into gel, doctor blading, pressing, and high-temperature calcination (ca. 1100 °C). For the formulation of mesoporous c-LLZO powder into gel first the powder is dispersed with the help of an organic solvent like ethanol and after drying is mixed with a binder (e.g. polyvinyl butyral) and a plasticizer (e.g. benzyl butyl phthalate). The formed porous ceramic scaffolds are thin of the order of 10 to 300 pm and preferentially between 100 and 200 highly suitable for solid-state battery fabrication.Synthesis of mesoporous LLZO powder-the effect of PEG to PG molar ratio

[0305] Through XRD and SEM analysis of the LLZO powders obtained under different synthesis procedures, the one obtained according to Fig. 24 with molar ratio of PEG (MW 20000): PG 2:0.5 consisted of pure cubic garnet phase (Fig. 25) and had a mesoporous structure (Figs.26A-26R and 27A-27F). Following hydrothermal aging, the first calcination was conducted at 300 °C for 2 h under an air atmosphere to burn out the organics and induce the formation of pores. Full crystallization of cubic phase LLZO was achieved after the second calcination conducted at 800 °C for 6 h. To evaluate the porous structure of the as-prepared cubic LLZO powders, BET surface area and BJH pore size analysis were employed. As shown in Fig. 27A, the N2adsorption / desorption isotherms curve is of the type IV characteristic of mesoporous structure. This is corroborated by the BJH curve in Fig. 27B and the SEM images in Figs. 27C-27F.Fabrication of thin porous LLZO scaffold: p-LLZO scaffold-2

[0306] The as-produced porous cubic LLZO powders are dispersed in ethanol (EtOH) containing 2 wt.% polyacrylic acid (MW = 2000 g / mol) as dispersant, using an ultrasonic horn at 100 W for 30 min. The suspension is left to settle for 4 h so the powder is recovered and dried. Collected powder is formulated afterwards into a homogeneous gel by mixing with, polyvinyl butyral (PVB), benzyl butyl phthalate (BBP), and acetone at 7:2:1 weight ratio in a ZrO2jar and ball-milled with 3.0 mm diameter spherical ZrO2beads for 1 h. The gel is doctor bladed on Mylar™ (polyester sheet) substrates and upon drying the resultant green films are manually peeled off the Mylar™ substrate, punched into selected sizes, and cold pressed under 10-ton load applied for 5 min, as illustrated in Figs. 28A, 28B, and 29. Punched green films were placed in between graphite foils and heated to 1100 °C under Ar flow (100 mL min-1). The graphite foil is used to maintain the scaffold integrity and avoid contamination or reaction with the ceramic crucible. The resulting films may have as consequence residual carbon deposits that are removed by reheating to 700 °C for 30 min under an air atmosphere. As shown in Fig. 30, the XRD results confirm that the p-LLZO scaffold-2 is a cubic garnet-type single-phase structure without other impurities showing up. The cubic-phase garnet crystal structure is targeted as it provides the highest Li+ ion conductivity. SEM images of the porous LLZO framework are shown in Figs. 31A-31 D along the digital picture of the p-LLZO-scaffold-2 in inset. The images clearly reveal the porous structure of ceramic films that are circa 130 pm, i.e. one order thinner than the p-LLZO sccaffold-1 counterparts. Their porosity was measured through the Archimedes method to be on average 45%. Finally, the EDS results in Figs. 32A-32D show that the elements are homogeneously distributed throughout the 3D ceramic network.Preparation of Hybrid Solid Electrolytes using p-LLZO Scaffold-2: HSE-2

[0307] In this experiment, the construction of hybrid polymer-in-ceramic solid electrolytes is described based on the thin porous-LLZO scaffold. Different polymeric or liquid electrolytes may be used to fill the porosity of the scaffold. In an embodiment the polymer used is poly(vinylidene difluoride) (PVDF) along a Li salt like lithium bis(trifluoromethane sulfonyl)imide (LiTFSI) dissolved prior in different solvents such as acetonitrile, dimethylsulfoxide (DMSO), dihydrolevoglucosenone (Gyrene), N-N’-dimethylpropyleneurea (DMPLI), triethyl phosphate (TEP), and y-valerolactone (GVL), or N-methyl pyrrolidone (NMP).

[0308] In a preferential embodiment, the hybrid solid electrolyte (HSE-2) is prepared as described in the schematic of Fig. 10. The as-prepared thin porous-LLZO scaffold-2 is immersed in a polymer solution of PVDF in which a Li-ion salt is added and let to infiltrate into the porous space of the scaffold before the solvent is evaporated to obtain HSE-2.Fabrication of hybrid solid electrolyte HSE-2 (p-LLZO scaffold-2 / PVDF-LiTFSI)

[0309] To prepare the HSE-2, various types of conductive polymer and Li salts combinations could be selected. Here we use PVDF-LiTFSI as an example to prove the concept. Poly(vinylidene difluoride) (PVDF) (MW = ~ 543000) and lithium bis(trifluoromethane sulfonyl)imide (LiTFSI) with a weight ratio of 3:1 are dissolved in N-methyl pyrrolidinone (NMP) with a polymer concentration of 15%. Then, the mixture is stirred for 24 h at 55 °C in a sealed bottle to obtain a homogeneous gel. Subsequently, the as-prepared p-LLZO scaffold-2 is immersed into PVDF / LiTFSI gel for 48 h and dried in the vacuum oven at 60 °C for 24 h to remove any residual trace of the NMP solvent.

[0310] Characterization of the constructed HSE-2 confirmed that the crystal structure of the cubic-garnet single-phase remains intact after PVDF infiltration (XRD patterns in Fig. 33). Meanwhile, the TGA analysis presented in Fig. 34 indicates that the weight percentage of polymer in HSE-2 is only 7%, which proves its predominant porous ceramic framework that not only yields high conductivity but also high mechanical strength for improved safety.

[0311] To clarify the chemical bonding between infiltrated PVDF-LiTFSI PSE and p-LLZO scaffold-2, XPS analysis was performed. As it can be seen in Figs. 35A and 35E, the C 1s spectra of PSE are characterized by peaks located at 284.78, 285.98, 288.68, and 290.38 eV which are attributed to (-CH3), (-CH2-CH2-), and (-CF3), (-CF2-CF2-). But in the case of HSE-2, the four peaks shift to lower binding energy by 0.2, 0.4, 0.2, and 0.4 eV respectively, as well as the intensity of peaks arising from -CH2-CH2- and -CF2-CF2- decreases while peaks arising from -CH3 and -CF3increases, indicating the formation of bonds between polymer groups and some metal atoms in p-LLZO- scaffold-2. To identify the type of bonds, the F 1s and N 1s spectra are presented in Figs. 35B, 35C, 35F, and 35G. The F 1s and N 1s spectra of PSE only exhibit one peak located at 687.68 eV and 399.88 eV, respectively, which come from the (-CF3) and (N-S) in TFSh groups. But in HSE-2, the F 1s and N 1s spectra show each a new peak appearing at 684.68 eV and 397.28 eV corresponding to La-F and La-N binding, respectively. This is clearly demonstrates the molecular interaction between the La atoms from p-LLZO scaffold-2 and PVDF-LiTFSI PSE. Specifically, the La-F bonds formed between La atoms and F from PVDF chains and not from TFSk groups, since in the present analysis of HSE-1 , with PEO-LiTFSI PSE and p-LLZO scaffold-1 , no La-F was detected. Moreover, Fig. 35D shows the La 3d spectra of the HSE-2 display two sets of peaks located at 855.18, 851.38 eV (La 3d3 / 2), and 838.48, 834.48 eV (La 3ds / 2), which shift to lower binding energy by 0.2, 0.1 , 0.2, and 0.1 eV compared with the peaks of p-LLZO scaffold-2, evidencing again the binding of the La atoms. Thus, it is proposed that robust La-F and La-N bonding forming between the PVDF PSE and p-LLZO scaffold-2 results in strong and homogeneous contact of PSE inside / outside the ceramic skeleton, as presented in Fig. 35H, hence superior hybrid solid electrolyte is obtained. The XPS surveys of the p-LLZO scaffold-2, PSE film, and HSE-2 are shown in Fig. 36. SEM images of the cross-section and top-surface of the HSE-2 and the SEM image of the cross-section of the commercial separator are shown in Figs. 37A-37D for comparison. The results show that the thickness of the HSE-2 is about 145 pm which is half of the thickness of the commercial separator.

[0312] The present ultra-thin HSE-2 demonstrates improvement in that it is a high-strength ceramic skeleton can resist the growth of Li dendrites as well as its reduced thickness can improve the energy density of the ASSLBs but also it could be used as a separator in standard LIB architecture. The EDS results shown in Figs. 38A-38F confirm the chemical homogeneity of HSE-2. Moreover, the ceramic and polymer phases are continuous, which are interconnected and interspersed in the as-prepared HSE-2. The continuous interface between p-LLZO scaffold-2 and PVDF-LiTFSI PSE can facilitate the fast migration of Li+due to the lower interfacial resistance, and the softer PSE can improve the interfacial connection with electrodes leading to high- performance ASSLBs.

[0313] Ionic conductivity is a key property of solid electrolytes. This property was confirmed for HSE-2 using electrochemical impedance spectra (EIS) and stainless-steel | HSE-2 | Stainless- steel cells. The results under various temperatures from 25 to 80 °C of three types of electrolytes: the HSE-2, pure cubic LLZO SE, and pure PVDF-LiTFSI PSE are shown in Fig. 39. Among them,the HSE-2 has the highest ionic conductivity, 0.437 mS cm-1at 25 °C and the lowest activation energy of 0.1905 eV. By comparison the respective values for pure cubic LLZO SE are 0.421 mS cm'1and 0.1997 eV. This good capability of Li+migration in HSE-2 is due to the molecular bonding of the conductive polymer within the porous-LLZO scaffold-2 and its low interface impedance.

[0314] Li-ion transference number (tLi+) is also a key property of the electrolyte. The lithium- ion transference number was measured using Li | HSE-2 | Li symmetric cells (see Fig. 40) by DC polarization and AC impedance. Thus HSE-2 exhibits a very high tLi+ of 0.72, which is much higher than the corresponding ones for PVDF-LiTFSI PSE (-0.25) or other hybrid electrolytes featuring cubic LLZO particles inserted in a polymer matrix (-0.31). This is mainly because the PSE does not have ion transport selectivity, and thus it usually has low tLi+. In contrast, a ceramic electrolyte as the c-LLZO compounded with polymer usually improves the tLi+ of the composite electrolyte. High tLi+ enables lower concentration polarization in practical application as well as it suppresses the driving force of Li dendrite growth.

[0315] To evaluate the interfacial stability and dendrite inhibition ability against lithium metal anode of the HSE-2, symmetric cells (Fig. 40) were employed. Long-term interfacial stability is critical for the stable operation of solid-state batteries, and this can be analyzed by constant current cycling tests as shown in Figs. 41 and 42. The results indicate that the HSE-2 did not short-circuit at current densities of 0.1 and 0.5 mA cm-2for 1000 h. Even at the higher current density of 0.5 mA cm-2, the symmetric cell still recovers to stable cycling after a short-term slight polarization.

[0316] Moreover, a wide electrochemical window (up to 4.9 V vs. Li+1 Li) of the HSE-2 was demonstrated through linear sweep voltammetry (LSV) as presented in Fig. 43. This excellent stability against Li metal is mainly due to the wide energy band system enabled by the strong interaction (bonding) of TFSI' groups and PVDF chains in infiltrated PSE with surface La atoms in the LLZO porous ceramic skeleton. The SEM images and EDS results of the interface (between the HSE-2 and Li metal) of the cycled cell at 0.5 mA cm'2(presented in Figs. 44A-44I) corroborate this conclusion as they show no signs of degradation like physical cracks or elemental transitioning. Therefore, the results clearly demonstrate that HSE-2 has excellent stability against Li metal, owing to fast ionic transport, low interfacial impedance, and good chemical / electrochemical stability. Moreover, high mechanical strength attributed to the ceramic skeleton and ultra-thin thickness are other merits that provide resistance to dendrite growth, high energy density due to high voltage, as well as safety.Preparation of the TiCk-coated NMC622 cathode

[0317] T etrabutyl titanate as the Ti source (0.1918 g corresponds to 1 .5 wt % TiCh) was added into 40 mL ethyl alcohol and stirred for 1 h. Then, 3 g LiNi0.6Co0.2Mn0.2O2 (NMC622) (MSE Supplies) powder was added into the mixed solution and continually stirred for another 2 h. Next, 0.4 mL ammonium hydroxide (NH4OH) was slowly added into the solution under stirring for another 2 h. Afterward, the residual powder was collected by centrifugation and washed with ethyl alcohol and DI water thoroughly, then dried under vacuum at 120 °C overnight. The final product was obtained by annealing the dried powder at 500 °C under an air atmosphere for 5 h. The obtained TiCh-coated NMC622 powder, carbon black (MSE Supplies), and PVDF (HSV900 PVDF binder, MTI Canada) with LiTFSI (weight ratio 3:1) were mixed in NMP solvent with a weight ratio of 8:1 :1 and coated on Al foil after stirring into a homogeneous suspension. Then, the electrode was dried at 50 °C for 6 h and further dried in a vacuum oven at 80 °C overnight to remove the NMP solvent. The dried electrode was punched into disks with a diameter of 9 mm, the mass loading of which was ~13 mg cm'2. The characterization of the TiC>2@NMC622 cathode (crystalline structure and particle morphology) is presented in Figs. 45A-45F.Characterization and performance of the TiC^-coated NMC622 cathode / HSE-2 / Li ASSBs

[0318] To demonstrate the applicability of the ceramic-based hybrid solid electrolyte (HSE-2) in practical high-voltage ASSBs, the TiCh-coated LiNi0.6Co0.2Mn0.2O2 (TiO2@NMC622) was used as the cathode and Li-metal as the anode to fabricate coin-type cells as per schematic presented in Fig. 46. The structural integrity of the as-assembled TiO2@NMC622 / HSE-2 / Li ASSB (before cycling) was evaluated with the SEM (Fig. 47A) and corresponding EDS mapping (Figs. 47B-47I), which showed a tight contact of the interfaces between HSE- and cathode or Li-metal. The interfacial boundaries are clear without elemental transitions before charge-discharge cycling. The morphologies and elemental distribution at the cathode / HSE-2 interface observed by SEM and EDS showed that the conductive PVDF-LiTFSI PSE from both cathode and HSE-2 components forms a homogeneous interfacial layer that provides integrity and unimpeded interfacial Li-ion conduction.

[0319] The charge-discharge cycling results of the as-assembled TiC>2@NMC622 / HSE-2 / Li ASSBs from 3 to 4.5 V at 0.1 C and 25 °C demonstrated that the capacity was 132.6 mAh g-1after 50 cycles, with a capacity retention of 82% and an average Coulombic efficiency of >99% as presented in Figs. 48A and 48B. The capacity fade is similar to that of the conventional NMCbatteries with organic liquid electrolytes (LE) as presented in Figs. 49A and 49B. To further explore the stability of the as-designed ASSLBs under high cycling rates, a rate performance test at 0.1 C, 0.2 C, 0.5 C, 1 C, and 2 C as presented in Figs. 50A and 50B. The specific discharge capacities were 161.36, 153.33, and 95.33 mAh g-1at 0.1 C, 0.2 C, and 0.5 C, respectively and remained stable, while the specific capacity dropped significantly at 1 C and 2 C. However, when the charge-discharge rate was restored to 0.1 C after experiencing the failure at 2 C chargedischarge, the specific discharge capacities can still recover to 157.57 mAh g-1, which indicated that the as-designed ASSLBs built with the HSE-2 have good reversibility and capacity.

[0320] To enhance the ionic conductivity of the PVDF-LiTFSI polymer electrolyte of the cathode and hence, the rate capability of the ASSLB made with the as-designed HSE-2 — the charge-discharge cycling of the as-assembled TiC>2@NMC622 / HSE-2 / Li ASSLBs was conducted at 40 °C in the voltage range from 3 and up to 4.5, 4.6, 4.7, and 4.8 V (vs. Li / Li+), respectively. The charge-discharge profiles (shown in Figs. 51A-51 D) of the full cell in the first 200 cycles at a rate of 0.2 °C were obtained. Stable reversible specific discharge capacities of 139.08, 139.80, 105.84, and 108.04 mAh g-1were achieved from each cutoff voltage (4.5-4.8 V) with a high Coulombic efficiency >99%. The corresponding energy densities at 100 and 200 cycles and 0.2 °C were 398.3 and 376.1 Wh kg-1respectively, which are higher than the current USABC goals for advanced high-performance batteries for electric vehicle applications (350 Wh kg-1). The cyclic voltammetry (CV) and accompanied EIS measurements before and after the CV tests for the as- assembled ASSLBs were conducted from 3 to 4.8 V at a scan rate of 0.1 mV s-1(shown in Figs. 52A-52C) demonstrating a highly reversible lithiation and de-lithiation functioning. The results demonstrated that the cycling stability of the ASSBs is remarkably improved in comparison with the batteries tested at room temperature (RT), even with the LE. The reversible specific discharge capacities dropped slightly during the first 40 cycles, a phenomenon attributed first to NMC622’s intrinsic property of surface degradation. At a highly delithiated state, the active Ni4+ions are unstable and have a tendency to form an NiO-like rocksalt phase on the cathode surface that increases the resistance, leading to capacity decay. The apparent morphological evolution of the TiCh-coated NMC622 particles before and after cycling can be observed in Figs. 45B-45F and Figs. 53A-53C respectively. Second, an in situ formed N and F-rich interphase was detected between cathode and HSE-2 in the ASSLBs, which is attributed to the decomposition of the infiltrated PVDF-LiTFSI during the initial cycles in analogy of the solid electrolyte interphase that forms in LE-based LIBs. After the initial 40 cycles, the reversible capacities remained stable at least up to 200 cycles. Notably, the average Coulombic efficiency of the ASSLBs tested from 3 to4.8 V reached 99.76% after 200 cycles, confirming the highly reversible and stable Li+intercalation / de-intercalation behavior. Furthermore, the rate performance of the as-assembled ASSLBs with various cutoff voltages (4.5-4.8 V) was also evaluated at 40 °C (Figs. 54A-54D). The specific discharge capacity decreases with increasing current density, but the ASSLBs tested at 40 °C provided higher discharge capacity when compared to that tested at RT (Fig. 50B).

[0321] Specifically, the as-assembled ASSLB tested from 3 to 4.8 V achieved reversible capacities of -52.72 and 46.34 mAh g-1at high discharge rates of 1 C and 2 C, respectively. The results illustrated that the as-assembled TiC>2@NMC622 / HSE-2 / Li ASSBs has excellent rate recovery. Therefore, the fast Li+transport and good interfacial stability with electrodes comes from the as-designed ceramic-based CSE that enables stable cycling ASSB performance at a high cutoff voltage.

[0322] With respect to the severe capacity decline exhibited by the ASSLB at 1 C and 2 C cycling, the Li / HSE-2 / TiC>2@NMC622 batteries were subjected to rate measurements under 25 °C RT (Fig. 50A and 50B) and 40 °C (Figs. 54A-54D). As shown in Figs. 54A-54D, at 40 °C, the specific discharge capacities were 94.76 mAh g_1at 1 C and 60.36 mAh g-1at 2 C, which are significantly improved compared to the nearly zero values obtained at 25 °C and the same cutoff voltage (4.5 V). Even when increasing the cutoff voltage to 4.8 V, the specific discharge capacities were 56.33 mAh g-1(1 C) and 39.64 mAh g-1(2 C). The reason for the drastic capacity decay with a rate increase was attributed to the cathode and not the HSE-2, is the limited conductivity of the PVDF-LiTFSI, which is the electrolyte used in the cathode. Specifically, Li-ion conductivity of PVDF-LiTFSI is one order of magnitude lower than that of HSE-2, plus it has a smaller Li-ion transfer number. Meanwhile, after the cycling rate was dropped (Figs. 53C-53D), both ASSLBs (operated at 25 °C and 40 °C) restored their capacities at near 100% Coulombic efficiency, proving the interfacial integrity / robustness of the designed HSE-2. Thus, the capacity loss under high rates can only be linked to the limited Li-ion transport capability of PVDF-LiTFSI and not the failure or degradation of the HSE-2. Further enhancement in rate performance should be attained with co-optimization of the conductivity of the cathode composite itself.

[0323] Stable SEs are essential for applications with high-voltage cathodes. Under such conditions, SEs may experience severe oxidation, decomposition, and deactivation during charging at high voltage, leading to inadequate cycling performance and even failure. The as- assembled TiC>2@NMC622 / HSE-2 / Li ASSBs exhibit excellent stability with high-voltage cathodes even under high cutoff voltage at 4.8 V. To determine the mechanism behind this high-voltagestability, the interface between the TiC>2@NMC622 cathode and the HSE-2 was studied. The cross-sectional SEM and the corresponding EDS images of the TiC>2@NMC622 / HSE-2 / Li ASSB after 200 cycles in the voltage range of 3-4.8 V at 0.2 C (40 °C) (Li foil was removed) is presented in Fig. 55A. Intimate contact between cathode and ceramic-based CSE was observed. In addition, the elemental mapping on Fig. 55A showed that Ni, Co, Mn, La, and Zr have created an intradiffused fine zone (see dotted lines in Fig. 55A). To further investigate the interface, the focused-ion beam (FIB) technique was employed to cut the smooth interface (Figs. 55B-55C). Figs. 55D-55F show the quantitative EDS spectra of three regions (cathode, interface, and HSE- 2) before and after cycling. Specifically, the concentration of Ni, Co, and Mn increases in the interface region of the cycled battery compared to the fresh one. Figs. 53A-53C show the cross- sectional SEM image of the ceramic-based HSE-2 / TiC>2@NMC622 interfacial area after cycling and details of the morphology and size (particle size ~9 mm) of the TiC>2@NMC622 active material. No delamination or cracking at the interfacial zone or particles was evident. Meanwhile, the corresponding EDS mappings of the FIB-cut surface exhibited an F- and N-rich interphase formed between cathode and HSE-2.

[0324] An XPS analysis of F1s, N1s, and Li 1s for the uncycled and cycled interface between cathode and HSE-2 was performed as presented in Figs. 56A-56H. Compared to the fresh / uncycled interface, the intensity of the peak at 684.95 eV in F1s spectra increased as a result of the formation of LiF in the interface, rendering it F rich; in the N1s spectra, a new peak appeared at 402.35 eV, which was ascribed to newly formed Li3N; the formation of these Li salts is manifested also in the Lils spectra with the appearance of a new peak at 55.22 eV corresponding to LiF and Li3N after long delithiation and lithiation cycling. The formation of LiF and Li3N at the interphase mainly comes from the decomposition of PVDF in cathode and PVDF- LiTFSI PSE in HSE-2 as evident by comparing C1s spectra before and after cycling, where the intensity of peak -CF3increases and peak -CF2-CF2- decreases. Thus, the N- and F-rich interphase (Li3N and LiF) formed at the interface between the TiO3@NMC622 cathode and HSE- 2 stabilizes the NMC particles against capacity decay under high cutoff voltage (4.8 V) and ensures active face-to-face contact of two components (cathode and HSE-2) during the long delithiation and lithiation cycling. Compared to traditional high-voltage Li-metal batteries that are stabilized by the introduction of LiF and Li3N interphase, the present battery has a LiF-Li3N-rich interphase between HSE-2 and electrodes via an in situ topotactic reaction during cycling. This in situ formed LiF-Li3N-rich interphase provides a superior stabilizing tool compared to traditional high voltage ASSB designs.Fabrication of Thin Porous LLZO Scaffold-3

[0325] For higher energy density of the all-solid-state batteries, the thinner solid-state electrolytes (SSEs) are desired. Accordingly, another variation of the above-mentioned fabrication method was designed to favour engineering of thin (< 90 pm scale) porous cubic LLZO scaffold- 3 (p-LLZO scaffold-3) was produced. The scaffold was fabricated starting from nanoscale cubic- LLZO powders obtained from a garnet-type solid electrolyte synthesis method as described in WO2023023856. The nanoscale c-LLZO powders were then formed into high-strength porous ceramic scaffolds (p-LLZO-scaffold) via formulation into gel, doctor blading, pressing, and fast high-temperature calcination (calcination at 1100 °C for 1 h). The present method produced thin ceramic membrane (< 90 pm) in contrast to traditional methods that produce thicker ceramic membranes.

[0326] The as-produced nanoscale cubic LLZO powders were dispersed in ethanol (EtOH) containing 2 wt.% polyacrylic acid (MW = 2000 g / mol) as dispersant, using an ultrasonic horn at 100 W for 30 min. The suspension was left to settle for 4 h so the powder was recovered and dried. Collected powder was formulated afterwards into a homogeneous gel by mixing with, polyvinyl butyral (PVB), benzyl butyl phthalate (BBP), and acetone at 7:2:1 weight ratio in a ZrC>2 jar and ball-milled (or other mixing treatments) with 3.0 mm diameter spherical ZrC>2 beads for 1 h. The gel was doctor bladed on Mylar™ (polyester sheet) substrates and upon drying the resultant green films were manually peeled off the Mylar™ substrate, punched into selected sizes, and cold pressed under 10-ton load applied for 5 min. Punched green films were heated to 1100 °C under O2 flow (100 mL min-1). The whole preparation is illustrated in Fig. 57. As shown in Fig. 58, the XRD results confirm that the p-LLZO scaffold-3 is a cubic garnet-type single-phase structure without other impurities showing up. The cubic-phase garnet crystal structure is targeted as it provides the highest Li+ion conductivity. SEM images of the porous LLZO scaffold-3 are shown in Figs. 59A-59F. The images clearly reveal the porous structure of ceramic films that are circa 89.5 pm, i.e. one order thinner than the porous LLZO scaffold-1 counterparts.Preparation of Hybrid Solid Electrolyte (HSE-3) using p-LLZO Scaffold-3

[0327] In this experiment, the construction of hybrid polymer-in-ceramic solid electrolytes is described based on the thin porous-LLZO scaffold. Different polymeric or liquid electrolytes may be used to fill the porosity of the scaffold. The polymer used can be, for example, poly(vinylidene fluoride) (PVDF) along a Li salt like lithium bis(fluorosulfonyl)imide (LiFSI) dissolved prior indifferent solvents such as acetonitrile, dimethylsulfoxide (DMSO), dihydrolevoglucosenone (Cyrene), N-N’-dimethylpropyleneurea (DMPLI), triethyl phosphate (TEP), and y-valerolactone (GVL), N-methyl pyrrolidone (NMP), or Dimethylformamide (DMF).

[0328] The as-prepared thin porous-LLZO scaffold in this example was immersed in a polymer solution of PVDF in which a Li-ion salt was added and allowed to infiltrate into the porous space of the scaffold before the solvent was evaporated to obtain the HSE-3.Fabrication of Hybrid Solid Electrolyte HSE-3 (PVDF-LiFSI-in-p-LLZO scaffold-3)

[0329] To prepare the HSE-3, various types of conductive polymer and Li salts combinations could be selected. Here we use PVDF-LiFSI as an example to prove the concept. Poly(vinylidene fluoride) (PVDF) (MW = ~ 543000) and lithium bis(fluorosulfonyl)imide (LiFSI) with a weight ratio of 3:1 are dissolved in Dimethylformamide (DMF) with a polymer concentration of 15%. Then, the mixture is stirred for 24 h at 55 °C in a sealed bottle to obtain a homogeneous gel. Subsequently, the as-prepared p-LLZO scaffold was immersed into PVDF / LiFSI gel for 48 h and dried in the vacuum oven at 60 °C for 24 h to remove any residual trace of the DMF solvent.

[0330] Characterization of the constructed HSE-3 confirmed that the crystal structure of the cubic-garnet single-phase remains intact after PVDF infiltration (XRD patterns in Fig. 60). Meanwhile, the TGA analysis presented in Fig. 61 indicates that the weight percentage of polymer in HSE-3 was only 4%, which proves its predominant porous ceramic framework that not only yields high conductivity but also high mechanical strength for improved safety. SEM images of the cross-section and top-surface of the HS3 and the SEM image of the cross-section of the commercial separator are shown in Figs. 62A-62F for comparison. The results show that the thickness of the HSE-3 is about 90.6 pm which is much thinner compared to commercial separator. The present ultra-thin HSE demonstrates improvement in that it is a high-strength ceramic skeleton that can resist the growth of Li dendrites as well as its reduced thickness can improve the energy density of the ASSLBs but also it could be used as a separator in standard LIB architecture. The EDS results shown in Figs. 63A-63F confirmed the chemical homogeneity of HSE-3. Moreover, the ceramic and polymer phases were continuous, which are interconnected and interspersed in the as-prepared HSE-3. The continuous interface between p-LLZO scaffold and PVDF-LiFSI PSE can facilitate the fast migration of Li+due to the lower interfacial resistance, and the softer PSE can improve the interfacial connection with electrodes leading to high- performance ASSLBs.

[0331] To clarify the chemical bonding in HSE-3 between infiltrated PVDF-LiFSI PSE and p- LLZO scaffold, XPS analysis was performed. As it can be seen in Fig. 64A, the C 1s spectra of PSE are characterized by peaks located at 284.38, 286.08, 288.38, and 290.48 eV which are attributed to (-CH3), (-CH2-CH2-), (-CF3), and (-CF2-CF2-) respectively. But in the case of HSE (Fig. 64B), the four peaks shifted to lower binding energy, as well as the intensity of peaks arising from -CH2-CH2- and -CF2-CF2- decreases while peaks arising from -CH3and -CF3increases, indicating the formation of bonds between polymer groups and some metal atoms in p-LLZO- scaffold. To identify the type of bonds, the F 1s and N 1s spectra are presented in Figs. 64C, 64D, 64E, and 64F. The F 1s and N 1s spectra of PSE only exhibit one peak located at 687.58 eV and 399.98 eV, respectively, which come from the (-CF3) and (N-S) in FSh groups. But in HSE the F 1 s and N 1 s spectra show each a new peak appearing at 684.58 eV and 397.18 eV corresponding to La-F and La-N binding, respectively. This is clearly demonstrating the molecular interaction between the La atoms from p-LLZO scaffold-3 and PVDF-LiFSI PSE. Specifically, the La-F bonds formed between La atoms and F from PVDF chains and not from FSL groups. Moreover, Fig. 64G shows the La 3d spectra of the HSE-3 display two sets of peaks located at 855.08, 851.28 eV (La 3d3 / 2), and 838.18, 834.38 eV (La 3ds / 2), which shifted to lower binding energy by 0.1 , 0.4, 0.3, and 0.5 eV compared with the peaks of p-LLZO scaffold-3, evidencing again the binding of the La atoms. Thus, it is proposed that robust La-F and La-N bonding forming between the PVDF-LiFSI PSE and p-LLZO scaffold-3 results in strong and homogeneous contact of PSE inside / outside the ceramic skeleton, hence superior hybrid solid electrolyte was obtained.

[0332] Ionic conductivity is a key property of solid electrolytes. This property was confirmed for HSE using electrochemical impedance spectra (EIS) and stainless-steel (SS) | HSE-3 | stainless-steel (SS) cells. The results under various temperatures from RT to 80 °C of pure PVDF- LiFSI PSE, and the HSE-3 (PVDF-LiFSI in porous LLZO scaffold) are shown in Figs. 65A and 65B. Among them, the HSE-3 has the highest ionic conductivity, 0.76 mS cm-1at 25 °C and the lowest activation energy of 0.22 eV (Fig. 65C). This good capability of Li+migration in HSE-3 is due to the molecular bonding of the conductive polymer within the porous-LLZO scaffold and its low interface impedance.

[0333] Li-ion transference number (tLi+) is also a key property of the electrolyte. The lithium- ion transference number was measured using Li | HSE-3 | Li and Li | PVDF-LiFSI | Li symmetric cells (see Figs. 66A to 66D) by DC polarization and AC impedance. Thus HSE-3 exhibited a very high tLi+ of 0.77, which is much higher than the corresponding ones for PVDF-LiFSI PSE (0.22). This is mainly because the PSE does not have ion transport selectivity, and thus it usually has\ow tLi+ . In contrast, a ceramic electrolyte as the c-LLZO compounded with polymer usually improves the tLi+ of the composite electrolyte. High tLi+ enables lower concentration polarization in practical application as well as it suppresses the driving force of Li dendrite growth.

[0334] To evaluate the interfacial stability and dendrite inhibition ability against lithium metal anode of the HSE-3, Li | HSE-3 | Li symmetric cells were employed. Long-term interfacial stability is important for the stable operation of solid-state batteries, and this can be analyzed by constant current cycling tests as shown in Fig. 67. The results indicate that the HSE-3 did not short-circuit at current densities of 0.1 , 0.2, 0.5 and 1 mA cm-2for over 1400 h. Even at the higher current density of 1 mA cm-2, the symmetric cell still recovers to stable cycling after a short-term slight polarization. The robust nature of the formed stable interface is further demonstrated with EIS plots of the symmetric cells before and after long-term plating-stripping cycling under 0.2, 0.5, and 1 mA / cm2, as shown in Fig. 68A. A suitable equivalent circuit with individual resistances and constant phase elements (CPEs) is employed to elucidate the Nyquist plot (after cycling at 1 mA / cm2for over 1400h) is shown in Fig. 68B to demonstrated the stable interface between HSE- 3 and Li metal. Further stepwise increase of the current density up to 1.6 mA / cm2, as shown in Fig. 69, showed no signs of short circuiting by the Li | HSE-3 | Li symmetric battery, which revealed a significantly enhanced critical current density (CCD) by the as-designed HSE (PVDF-LiFSI in porous LLZO scaffold).

[0335] Moreover, a wide electrochemical window (up to 5.12 V vs. Li+1 Li) of the HSE-3 was demonstrated through linear sweep voltammetry (LSV) as presented in Fig. 70. This excellent stability against Li metal is mainly due to the wide energy band system enabled by the strong interaction (bonding) of FSk groups and PVDF chains in infiltrated PSE with surface La atoms in the LLZO porous scaffold-3. Therefore, the results clearly demonstrate that HSE-3 has excellent stability against Li metal, owing to fast ionic transport, low interfacial impedance, and good chemical / electrochemical stability. Moreover, high mechanical strength attributed to the ceramic skeleton and ultra-thin thickness are other merits that provide resistance to dendrite growth, high energy density due to high voltage, as well as safety.Preparation of Composite Cathode Materials

[0336] Low ionic conductivities and apparent chemical diffusion coefficients of Li+ion is one of the challenges for cathode materials in all-solid-state batteries. Here, a core-shell structural composite cathode material, taking LiFePCL (LFP) (core) and nanoscale LLZO (shell) as anexample, was prepared with a simple mechanical synthesis method (as shown in Fig. 71), exhibiting enhanced apparent chemical diffusion coefficientcompared to the commercial LFP material. Specifically, LFP: Nano LLZO: Carbon black = 4:1 :1 wt% were ball-milled at 650 rpm for 1 h, then the collected powders were sintered at 350 °C for 1 h. In Fig. 72, the XRD patterns of commercial LFP and nanoscale LLZO employed are presented. As for the composite cathode powders, the XRD results of them exhibit mixing characteristic peaks of LFP and cubic LLZO without impurities.Kinetic Analysis of LLZO-LFP Composite Cathode Material

[0337] To verify the enhanced diffusion coefficient of the as-prepared LLZO-LFP cathode material, cyclic voltammetry (CV) measurements were employed. The peak current ( / p) exhibits a linear relation with the square root of scanning rate (v1 / 2). This is a typical diffusion controlling response, which indicates that the kinetics of two-phase transition in LiFePO4 behave similar as a diffusion process. For a homogeneous system, the chemical diffusion coefficient can be calculated from the following equation, the Randles-Sevcik equation:Ip= 0.4463n3 / 2F3 / 2CSR1 / 2T1 / 2D1 / 2v1 / 2where: Ip. the peak current (A); n: the charge transfer number; F: the Faraday’s constant (96486 C mol'1); C: the Li-ion concentration in LiFePCL; S: the surface area of the electrode (cm2); F: the gas constant (8.314 J moHK'1); T: the absolute temperature (K); £>: the chemical (Li-ion) diffusion coefficient (cm2s'1); v. the scanning rate (V s'1). As shown in Figs. 71A-71 D, based on the CV curves of commercial LFP under the scanning rate range from 0.01 to 1 mV / s (Fig. 73A) and the relation between peak current vs. scanning rate (Fig. 73B), the apparent chemical diffusion coefficients of commercial LFP were calculated as 3.89 x 10'14(Cathodic) and 3.28 x 10'14(Anodic). Based on Figs. 73C and 73D, the apparent chemical diffusion coefficients of as- prepared LLZO-LFP composite cathode were calculated as 8.28 x 10'14(Cathodic) and 6.95 x 10'14(Anodic), exhibiting a over twofold increase.Composition Analysis of LLZO-LFP Composite Cathode Material

[0338] To understand the mechanism behind enhanced kinetic property of the composite LLZO-LFP cathode, XPS was employed to analyze the chemical composition of the composite cathode powders collected after ball-mill and sintering. As shown in Figs. 74A and 74B, there are new peaks present in Fe 2p and La 3d spectra demonstrating elemental diffusion or chemical interactions occurred in the composite LLZO-LFP cathode after ball-mill and sintering treatment.There are no similar elemental diffusion or chemical interactions formed in other spectra as shown in Figs. 74C-74F.

[0339] To further observe the elemental diffusion and chemical interactions in composite LLZO-LFP cathode powders after ball-mill and sintering treatment, TEM was employed. As shown in Fig. 75, the TEM image of the composite LLZO-LFP cathode particle exhibiting cubic LLZO phase in the shell regions (as marked in the top three boxes) and LFP phase in the core region (as marked in bottom box).Fabrication of All-Solid-State Battery (ASSB) Enabled by HSE-3 and Composite Cathode

[0340] To demonstrate the feasibility of the HSE-3 and as-prepared composite LLZO-LFP cathode material, composite cathode slurry was prepared with PVDF and NMP and then infiltrated into as-prepared porous LLZO scaffold-3 then the whole composite cathode infiltrated porous LLZO scaffold was filled with PVDF-LiFSI gel as presented in Fig. 76. The cross-sectional SEM images of the porous LLZO scaffold-3 after infiltration of composite cathode and PVDF-LiFSI are shown in Figs. 77A-77B, exhibiting the thickness of the composite cathode layer and HSE layer are ~35 pm and ~55 pm, respectively. The corresponding EDS mapping and linear scanning of La and Fe are presented in Figs. 77C-77E, demonstrating the apparent boundary between composite cathode layer and HSE-3 layer. Then the ASSBs were assembled as presented in Fig. 78, the Li metal was employed as anode. The assembled ASSBs were stored in vacuum oven under 60 °C overnight then cooled down naturally, to obtain good connection between HSE-3 layer with Li metal and composite cathode layer with Al foil (current collector). The as-assembled ASSBs were cycled at 60 °C, from 2.5 to 4 V under 0.2 C for long-term charge-discharge cycling. As shown in Figs. 79A-79B, the ASSBs enabled by composite LLZO-LFP cathode and HSE-3 cycled stably for over 150 cycles with high discharge capacity (-168 mAh / g) which is closed to the theoretical capacity of the commercial LFP. As for the rate performance, the assembled ASSBs were cycled at 60 °C, at various charge-discharge rates ranging from 0.1 C to 2 C. As shown in Figs. 80A-80B, even at high charge-discharge rate (2C), high specific discharge capacities - 110 mAh / g were achieved, which demonstrated the excellent kinetic property and functionality of the HSE-3 / composite LLZO-LFP cathode system.

Claims

WHAT IS CLAIMED IS:

1. A porous Li-ion conducting garnet ceramic having a porosity of from 20% to 45%, wherein the garnet has a formula Liy-xDyLa3Zr20i2 where x and y are 0 < x < 3 and 0 < y < 1 respectively and D is a dopant selected from Ta, Nb, Al, Sn, Ge, Si, Li, Na, and K.

2. The porous Li-ion conducting garnet ceramic of claim 1 , wherein the dopant is Al.

3. The porous Li-ion conducting garnet ceramic of claim 2, wherein the garnet has the formula Li6.iAlo.3La3Zr20i2.

4. The porous Li-ion conducting garnet ceramic of any one of claims 1 to 3, wherein the porosity is measured by the Archimedes method.

5. The porous Li-ion conducting garnet ceramic of any one of claims 1 to 4, wherein the porous Li-ion conducting garnet ceramic has a thickness of less than 2 mm.

6. The porous Li-ion conducting garnet ceramic of claim 5, wherein the thickness is less than 100 pm.

7. A composite material comprising the porous Li-ion conducting garnet ceramic as defined in any one of claims 1 to 6, and a polymer electrolyte doped with a Li salt that is infiltrated in the porosity of the porous Li-ion conducting garnet ceramic.

8. The composite material of claim 7, wherein the polymer electrolyte coats at least one surface of the porous Li-ion conducting garnet ceramic.

9. The composite material of claim 7 or 8, wherein the polymer electrolyte coats the porous Li-ion conducting garnet ceramic.

10. The composite material of any one of claims 7 to 9, wherein the polymer electrolyte is selected from polyethylene oxide, polyvinyldene difluoride, polyvinyl alcohol, polycaprolactone, polychitosan, polyvinyl pyrrolidone, polyvinyl chloride, and / or polyimide.

11. The composite material of any one of claims 7 to 10, wherein the polymer electrolyte is polyethylene oxide and / or polyvinylidene difluoride.

12. The composite of any one of claims 7 to 11 , wherein the polymer electrolyte is present in a concentration of from 1 wt. % to 20 wt. %.

13. An all-solid-state lithium battery comprising an anode, a cathode and the composite material as defined in any one of claims 7 to 12 at the interface between the cathode and the anode.

14. The all-solid-state lithium battery of claim 13, wherein the composite material is provided as a separator in the battery and as an electrolyte.

15. An electrolyte-cathode composite, comprising a bilayer wherein the composite material as defined in any one of claims 7 to 12 forms a first layer, and a cathode material forms a second layer, wherein the cathode material covers the first layer and infiltrates the porosity of the composite material of the first layer.

16. The electrolyte-cathode composite of claim 15, wherein the cathode material is lithium iron phosphate (LFP) or lithium nickel-manganese-cobalt oxide (NMC).

17. A method for producing the porous Li-ion conducting garnet ceramic as a powder as defined in any one of claims 1 to 6, comprising: providing an aqueous phase comprising Zr, La, and optionally a dopant; mixing in the aqueous phase an organic polymer with pore forming properties to obtain a mixture; hydrothermally aging the mixture to obtain an intermediate comprising crystalline La(OH)s and an amorphous Zr hydroxide; calcining the intermediate with Li to obtain the porous Li-ion conducting garnet ceramic.

18. A method for producing the porous Li-ion conducting garnet ceramic as a scaffold as defined in any one of claims 1 to 6, comprising: performing the method as defined in claim 17 to obtain a powder of the porous Li- ion conducting garnet ceramic;mixing the powder with the organic polymer with pore forming properties to obtain a powder mixture; depositing the powder mixture onto a sacrificial substrate; and calcining the porous scaffold to remove the filler polymers and obtain the porous Li-ion conducting garnet ceramic as said scaffold.

19. A method for producing the porous Li-ion conducting garnet ceramic as a scaffold as defined in any one of claims 1 to 6, comprising: mixing a nanoscale cubic LLZO with an organic polymer to obtain a mixture; depositing the mixture onto a sacrificial substrate; and calcining the mixture to remove the organic polymer and produce a porous Li-ion conducting garnet ceramic as a scaffold.

20. The method of claim 17, wherein the aqueous phase has a pH of from more than 7 to up to 14.

21. The method of any one of claims 17 to 20, wherein the calcining is performed at a temperature of from 1000 to 1100 °C.

22. The method of any one of claims 16 to 18, wherein the organic polymer is selected from polyethylene glycol, propylene glycol, polyvinyl butyral, benzyl butyl phthalate and combinations thereof.

23. The method of any one of claims 17 to 21 , wherein the depositing is a doctor-blading or tape-casting.

24. A method of producing the composite material as defined in any one of claims 7 to 12, comprising providing the porous Li-ion conducting garnet ceramic as defined in any one of claims 1 to 6, and infiltrating the polymer electrolyte in the porosity of the porous Li- ion conducting garnet ceramic.

25. A method of producing the composite material as defined in any one of claims 7 to 12, comprising producing the porous Li-ion conducting garnet ceramic as defined in anyone of claims 1 to 7 with the method as defined in any one of claims 17 to 23, and infiltrating the polymer electrolyte in the porosity of the porous Li-ion conducting garnet ceramic.

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