A composite coating material and its preparation method and application

Abstract

The present invention relates to a composite coating material, a preparation method thereof, and an application thereof. The composite coating material comprises: a core having a core-shell structure; the shell structure of the core being made of carbon, and the core structure of the core being made of a silicon-based material; and a shell coated outside the core, the shell comprising a pore structure containing a polymer and hydroxylated lithium lanthanum zirconium-oxygen particles; the pore structure comprising macropores with a pore diameter of 200-1000nm, mesopores with a pore diameter of 50-100nm, and micropores with a pore diameter of 2-5nm. The present invention significantly improves the electrochemical performance and mechanical stability of silicon-carbon negative electrode materials through composite coating structure design and interface coordinated regulation, and achieves dual adaptation of liquid and solid-state battery systems.

Description

Technical Field

[0001] The present invention belongs to the field of materials, and in particular relates to a composite coating material and a preparation method and application thereof. Background Art

[0002] Silicon-based materials are considered to be the most promising negative electrode systems due to their theoretical specific capacity of up to 4200 mAh / g. However, the electrode pulverization problem caused by their drastic volume expansion (>300%) seriously restricts their practical application. Although traditional carbon coating technology can partially alleviate the volume effect, a single carbon layer coating is difficult to meet the synergistic needs of ion transport and electron conduction, especially in fast charging scenarios. In existing technologies, tantalum-doped LLZO (Li 7-x La3Zr 2-y Ta y O 12 )Poor interface compatibility of the coating layer leads to an interface impedance of >180 Ω·cm 2 .

[0003] In the liquid electrolyte system, the existing technology uses electrospinning to prepare the core-shell structure of PAN-coated silicon-based materials. Although the physical coating of silicon particles is achieved, the ionic conductivity of the PAN-based coating layer is only 10 -6 S / cm2 level, unable to meet the demands of high-rate charge and discharge. More seriously, this solution's efficiency in suppressing silicon volume expansion is less than 50%, causing the capacity retention rate to decay to 72.3% after 500 cycles, making it difficult to coordinate mechanical buffering capacity with electrochemical stability.

[0004] Prior art CN112467116A discloses a graphite-coated material, its preparation method, and a battery negative electrode. The graphite-coated material comprises: an inner core comprising graphite; and an outer shell coating the inner core, comprising tantalum-doped lithium lanthanum zirconium oxide, amorphous carbon, and a conductive material. Doping the graphite surface coating with tantalum-doped lithium lanthanum zirconium oxide and a conductive material effectively improves the lithium ion transmission rate and diffusion coefficient, and the conductive material effectively improves the material's electronic conductivity. The tantalum-doped lithium lanthanum zirconium oxide, amorphous carbon, and conductive material in the outer shell synergistically exhibit excellent lithium ion and electronic conductivity, which helps improve the ion transmission rate and conductivity of the graphite-coated material, effectively enhancing the rate capability, safety, and cycle performance of the graphite negative electrode material. However, this existing technology is only applicable to graphite cores (theoretical capacity 372 mAh / g) and cannot be adapted to high-capacity silicon-based materials (theoretical capacity 4200 mAh / g). Furthermore, its expansion suppression is insufficient: it does not involve a volume expansion buffering mechanism and has no effect on the expansion of silicon-based materials greater than 300%. The interface impedance is high: unmodified LLZO has poor compatibility with the graphite interface, and the interface impedance of solid-state batteries is greater than 180 Ω·cm. 2 .

[0005] Prior art CN111244410B discloses a lithium battery negative electrode material and its preparation method, specifically a silicon-based negative electrode material with high stability and long cycle life and its preparation method. The silicon-based negative electrode material has a core-shell structure, the core contains silicon oxides and has a high initial efficiency; the outer shell consists of three layers, the inner layer is a carbon coating layer, which can effectively buffer the volume expansion of the core and improve electronic conductivity, and the middle layer is an HF isolation layer, which can effectively prevent F - Pass without obstructing Li + The outermost layer is an artificial SEI film with Li + The conductive layer can effectively improve lithium ion conductivity and stabilize the SEI film. Using this negative electrode material can produce lithium-ion batteries with high energy density and long life. However, this existing technology involves a complex multilayer deposition process, insufficient SEI layer ion conductivity, and easy delamination at the interface between the SEI layer and the carbon layer. The capacity retention rate after 500 cycles is only 73.2%.

[0006] Prior art CN115241526A discloses an artificial graphite composite material coated with a boron-doped solid electrolyte composite material, which presents a core-shell structure, wherein the core is artificial graphite and the shell is composed of 0.5-2% lithium borate, 1-5% solid electrolyte, 0.5-2% conductive agent and amorphous carbon. The composite material of this invention utilizes the high ionic conductivity of the solid electrolyte and the electronic conductivity of the conductive agent to improve the electronic ion conduction rate of the material. At the same time, lithium borate provides sufficient lithium ions for the material to improve its initial efficiency, and exerts its coupling effect as a coupling agent in the coating layer to form a conductive network, thereby improving the structural stability of the material and its rate cycling performance. However, this prior art is designed for graphite, and lithium borate (0.5-2wt%) cannot adapt to the expansion stress of silicon-based materials; the conductive network has defects, and the electron mobility of the conductive network constructed by the coupling agent (aluminate / titanate) is only 10 -3 S / cm order of magnitude.

[0007] For solid-state battery systems, PVDF-HFP-based solid electrolytes have good interface flexibility, but their room temperature ionic conductivity is less than 10 -5 S / cm, and the interface impedance with the silicon negative electrode exceeds 200 Ω·cm ² , making efficient charge transfer difficult to achieve. A deeper technical bottleneck lies in the significant interfacial incompatibility between traditional inorganic fillers (such as LLZO and LATP) and the polymer matrix. This leads to filler agglomeration (particle size > 500 nm) and stress concentration, severely weakening the mechanical integrity of the composite electrolyte. Furthermore, this organic solid electrolyte system suffers from poor interfacial contact with silicon-based anode materials, resulting in poor electrochemical performance of solid-state batteries.

[0008] Common technical challenges are concentrated in two areas: First, existing coating structures cannot establish continuous dual ion / electron transport channels, making it difficult to simultaneously improve ionic conductivity and electron mobility; second, weak interfacial interactions between the inorganic and organic phases lead to interfacial delamination during cycling, causing contact failure between the active material and the current collector. These deficiencies collectively restrict the practical application of silicon-based anodes in high-performance batteries, and breakthroughs are urgently needed through collaborative innovation in materials design and interface engineering. Summary of the Invention

[0009] The purpose of the present invention is to provide a composite coating material with high ionic conductivity and low volume expansion rate, a preparation method thereof, and an application thereof in silicon-carbon negative electrode materials.

[0010] In order to achieve the above object, the present invention adopts the following technical solutions:

[0011] A composite coating material, comprising:

[0012] The core is a core-shell structure; the shell structure of the core is made of carbon, and the core structure of the core is made of silicon-based material;

[0013] and a shell coated on the outer surface of the core, wherein the shell comprises a pore structure and contains a polymer and hydroxylated lithium lanthanum zirconium oxide particles;

[0014] Among them, the structural formula of lithium lanthanum zirconium oxide particles is Li 7-x La3Zr 2-y M y O 12 , M is at least one of the doping elements Ta, Al, and Nb that replaces Zr in equal amounts, 0≤x≤1, 0≤y≤0.6;

[0015] The pore structure includes macropores with a pore diameter of 200-1000 nm, mesopores with a pore diameter of 50-100 nm, and micropores with a pore diameter of 2-5 nm.

[0016] Within the pore structure, macropores (200-1000 nm) preferentially absorb the macroscopic strain of silicon core expansion; mesopores (50-100 nm) disperse shear stress, preventing crack propagation; and micropores (2-5 nm) stabilize the SEI membrane through surface adsorption. In Example 1, the three pore structures synergistically resulted in an electrode integrity rate of 92% after cycling, compared to only 78% for a dual-pore system.

[0017] The necessity and size impact of different apertures are as follows:

[0018] Macropores: Absorb the macroscopic strain of silicon expansion (ΔV>200%). They need to match the 5-12μm size of the silicon particles in the core. Too small a size will result in reduced buffering efficiency.

[0019] Mesopores: promote electrolyte penetration (liquid system) and ion transport (solid system). The size is adapted to the LLZO particle size of 80-90 nm. Too large will lead to a decrease in active sites, while too small will block the ion migration path.

[0020] Micropores: provide high specific surface active sites (Li + adsorption), the size is suitable for the graphitized carbon layer. Too large will lead to too thick or unstable SEI film, and too small will lead to Li + Diffusion is hindered.

[0021] The effects of the number and type of pores are as follows:

[0022] There are only two scales of pores:

[0023] a) Loss of macropores: volume expansion rate rises to >200%;

[0024] b) Missing micropores: 1C capacity decay rate increases by 30% (ion transport bottleneck).

[0025] Excessive pore sizes (e.g., quaternary pore structure):

[0026] a) Mechanical strength decreases and electrode processing breakage rate increases;

[0027] b) Improved interfacial impedance dispersion.

[0028] In one preferred embodiment, in the pore structure, the macroscopic pores account for 50-60% (volume ratio). In the embodiment, focused ion beam-scanning electron microscopy (FIB-SEM) three-dimensional reconstruction statistics are used to extract the pore volume by threshold segmentation.

[0029] In one of the preferred embodiments, in the pore structure, the mesopores account for 30-40% (volume ratio). In the embodiment, based on the BJH model calculation of the nitrogen adsorption-desorption isotherm, LLZO nanoparticles (D50=80-90 nm) are accumulated to form "spherical dense packing" pores (porosity ≈26%), and the spatial distribution uniformity is verified by electron tomography (ET).

[0030] In one preferred embodiment, in the pore structure, the micropores account for 5-10% (volume ratio). In the embodiment, small-angle X-ray scattering (SAXS) combined with density functional theory (DFT) model fitting is used to quantify the lattice-level pores caused by carbonization shrinkage (the micropore density is saturated when the graphitization degree of the carbon layer is ≥80%).

[0031] In one preferred embodiment, the macropores are mainly distributed on the outer surface of the coating layer; the mesopores are distributed in the LLZO particle accumulation area; and the micropores are concentrated in the carbonized polymer matrix.

[0032] The pore distribution basically conforms to the rule of sparse outside and dense inside.

[0033] In one preferred embodiment, the macroscopic pore density increases from the inside to the outside.

[0034] In one preferred embodiment, the polymer is one or more of polyacrylonitrile (PAN), polyvinylidene fluoride-hexafluoropropylene copolymer (PVDF-HFP), and polyethylene oxide (PEO).

[0035] In one preferred embodiment, in the inner core, the mass proportion of the core structure is 15-50wt%.

[0036] In one preferred embodiment, the particle size D50 of the core is 5-12 μm, and the specific surface area is 1-3 m 2 / g, and the tap density is 0.8-1.2 g / cm 3 .

[0037] In one preferred embodiment, the silicon-based material includes elemental silicon, porous silicon, nano silicon, SiO x A combination of one or more silicon-oxygen materials (0<x<2).

[0038] In one preferred embodiment, the shell structure has a thickness of 50-200 nm and a crystallinity of ≥75%.

[0039] In one preferred embodiment, the shell structure is made of graphitized carbon.

[0040] In one preferred embodiment, the porosity of the inner core is ≤5%.

[0041] In one preferred embodiment, the core is G4S-C800 or G4S-C600, purchased from Shanghai Shanshan New Materials Co., Ltd.

[0042] In one preferred embodiment, the mass ratio of the polymer to the hydroxylated lithium lanthanum zirconium oxy-particles is 8-5:2-5.

[0043] In the present invention, LLZO nanoparticles (Li 7-x La3Zr 2-y M y O 12 ) provides a fast migration path for lithium ions, and the polymer carbonization forms a continuous conductive network, realizing the decoupling and synergy of ion / electron transport, which is better than the traditional single carbon coating (10 -6S / cm) increased by 4 orders of magnitude. The formed "macropore (200-1200 nm)-mesopore (50-100 nm)-micropore (5-15 nm)" structure dissipates stress in stages through the synergistic effect of multi-scale pores, avoiding cracking, powdering and failure of the coating layer during long-term cycles. LLZO surface hydroxylation (≥3 / nm 2 ) forms La-OC covalent bonds with the carbonized polymer network, strengthening the inorganic / organic interface. The high melting point of LLZO (>1500°C) synergistically with the flame retardancy of the carbonized polymer inhibits thermal chain reactions; its porous structure slows heat accumulation. This results in a thermal runaway onset temperature of 200°C, a 56% increase compared to conventional cores (128°C).

[0044] In one preferred embodiment, the preparation method of hydroxylated lithium lanthanum zirconyl oxide particles comprises the following steps: ball milling and drying the lithium lanthanum zirconyl oxide particles, and then adding them to an acid solution for ultrasonic treatment to obtain hydroxylated lithium lanthanum zirconyl oxide particles.

[0045] In one preferred embodiment, the ball milling is performed using a gradient dispersion process, wherein the gradient dispersion process is: coarse grinding at a rotation speed of 300-400 rpm for 1-3 hours, and then fine grinding at a rotation speed of 100-150 rpm for 3-5 hours.

[0046] In one preferred embodiment, the ball milling medium used in the ball milling is zirconia balls, and the diameter of the zirconia balls is 2-6 mm.

[0047] In one preferred embodiment, in the zirconia balls, the volume ratio of the zirconia balls with a diameter of 2-3 mm to the zirconia balls with a diameter of 5-6 mm is 1:2-3.

[0048] In one preferred embodiment, the particle size D50 of the hydroxylated lithium lanthanum zirconyl oxide particles is 80-90 nm.

[0049] In one preferred embodiment, the acid solution is a nitric acid solution, and the concentration of the nitric acid solution is 1-5 wt %.

[0050] In one preferred embodiment, the temperature of ultrasonic treatment is 40-90° C., preferably 70-80° C.; and the ultrasonic treatment time is 30-180 min, preferably 60-120 min.

[0051] In one preferred embodiment, the frequency of ultrasonic treatment is 40-60 kHz, the power is 300-350 W, and the pulse mode is 5-10 s on / 5-10 s off.

[0052] In one preferred embodiment, the drying step is: treating at 60-110° C. for 6-24 hours.

[0053] In one preferred embodiment, the surface of the hydroxylated lithium lanthanum zirconium oxide particles has a hydroxyl density greater than 3 / nm 2 active site.

[0054] When the lithium lanthanum zirconyl particles are treated with nitric acid, the surface undergoes selective etching (La-O bond hydrolysis). Simultaneously, some Zr-O bonds break, forming surface hydroxyl groups (-OH). The hydroxylated lithium lanthanum zirconyl particles (La-OH on the surface) undergo a condensation reaction with polymers (such as PAN) during carbonization, forming La-OC bonds. These La-OC bonds increase the interfacial binding energy greater than that of physical adsorption in the unhydroxylated state.

[0055] (1) Quantitative relationship between hydroxyl density and interfacial binding energy

[0056] Assume that the hydroxyl density on the surface of LLZO is ρ OH (Unit: pieces / nm 2 ), the binding energy of each hydroxyl group to form a La-OC bond with the polymer carbonization network is εbond (needs to be experimentally determined, unit eV / bond), and the effective contact area is (AFM measurement is required, nm 2 ). Then the interface binding energy per unit area E can be expressed as:

[0057] E=ρOH*εbond*A contact ;

[0058] Assume ρOH≥3 / nm², εbond=0.8eV, A contact =80%;

[0059] E≥3*0.8*0.8=1.92eV / nm 2 =8.7J / m 2 .

[0060] (2) The promoting effect of La-OC bond on ion transport

[0061] The polar effect of La-OC bond can reduce the Li + Migration activation energy E. According to the Arrhenius equation, decreasing E can increase ionic conductivity.

[0062] (3) Interface stability inhibits volume expansion

[0063] The breaking strength of the La-OC bond is much greater than the volume expansion stress of silicon. La-OC at the silicon interface can effectively inhibit the volume expansion of silicon.

[0064] In a preferred embodiment, the polymer is uniformly distributed in the shell.

[0065] In one preferred embodiment, the hydroxylated lithium lanthanum zirconyl particles are uniformly distributed in the shell.

[0066] Based on the same inventive concept, the present invention also claims protection for a method for preparing the composite coating material, comprising the following steps:

[0067] S1. Mixing a polymer and hydroxylated lithium lanthanum zirconium oxide particles in a mass ratio of 8-5:2-5; adding a porogen after mixing evenly, and sonicating to obtain a mixture;

[0068] S2. adding the core to the mixture and mixing uniformly to obtain a pre-product;

[0069] S3, performing gradient heat treatment on the pre-product to obtain a composite coating material;

[0070] The gradient heat treatment process is as follows: under argon protection, heating to 200-250°C at a rate of 2-3°C / min for pre-oxidation for 1-2 hours, then heating to 700-800°C at a rate of 5-10°C / min for carbonization for 2-3 hours, then cooling to 600-650°C at a rate of 5-7°C / min for insulation for 15-20 minutes, and then slowly cooling to room temperature.

[0071] During the gradient heat treatment process, under argon protection, the pre-oxidation of the polymer molecular chain is completed at 200-250°C; then, at the main carbonization temperature of 700-800°C, constant temperature carbonization is carried out at a certain argon flow rate to form a continuous conductive network; finally, a gradient slow cooling program is performed to optimize the carbon crystal structure.

[0072] In one preferred embodiment, in step S1, the amount of the porogen added is 0.1-15 wt % of the polymer mass.

[0073] In one preferred embodiment, in step S1, the porogen is one or a combination of polyethylene glycol and polyvinyl pyrrolidone.

[0074] In one preferred embodiment, in step S1, the porogen is one or more of PEG200, PEG300, PEG400, PEG500, PEG600, PEG800, PEG1000, PEG1500, PEG2000, PEG3000, PVP K13, PVP K16, PVP K18 and PVP K30.

[0075] In one preferred embodiment, in step S1, the molecular weight of the porogen and the macroscopic pore size satisfy the relationship:

[0076] .

[0077] in, is the average diameter of macropores (nm); is the number average molecular weight of the porogen.

[0078] In one preferred embodiment, in step S1, the ultrasonic process is: power of 300-400 W, frequency of 40-50 kHz, and time of 15-60 min.

[0079] In one preferred embodiment, in step S1, before the polymer is mixed with the treated lithium lanthanum zirconium oxide particles, the polymer is first dispersed in a solvent to obtain a solution with a concentration of 5-15 wt%.

[0080] In one preferred embodiment, the solvent includes one or more of dimethylformamide (DMF), N-methylpyrrolidone (NMP), tetrahydrofuran (THF) and acetone.

[0081] In one preferred embodiment, the solvent includes acetone and one or more of dimethylformamide (DMF), N-methylpyrrolidone (NMP), and tetrahydrofuran (THF).

[0082] In a preferred embodiment, the volume concentration of acetone in the solvent is 50%-70%.

[0083] In one preferred embodiment, in step S2, the mass ratio of the core to the polymer in the mixture is 10:0.5-10:3.

[0084] In one preferred embodiment, in step S3, during the gradient heat treatment, the argon flow rate and the temperature satisfy the relationship: V(Ar)=0.5×T / 100 (L / min), where T is the real-time temperature (°C).

[0085] In one preferred embodiment, in step S3, the slow cooling stage is performed at a rate of 10-15°C / min to cool to room temperature.

[0086] During the gradient heat treatment, the precise construction of macropores is achieved by selecting a specific porogen (the relationship curve between the molecular weight of the porogen and the macropore size is shown in Figure 7 ); Combined with the gradient ball milling process (dynamic adjustment from 300 to 150 rpm), the median particle size (D50) of LLZO particles is controlled at 80 nm, and the 50-100 nm mesopore distribution is optimized; finally, through the critical phase transition temperature control at 650°C, synergistically with the microporous structure (5-15nm), a balance between ion / electron dual channels and mechanical stability is achieved.

[0087] Based on the same inventive concept, the present invention also claims protection for an electrode slurry, which includes the composite coating material.

[0088] In one preferred embodiment, the electrode slurry includes the composite coating material, conductive carbon black, polyvinylidene fluoride and N-methylpyrrolidone; the mass ratio of the composite coating material, conductive carbon black and polyvinylidene fluoride is 6-8:1-3:1-2.

[0089] N-methylpyrrolidone is used as a solvent, and composite coating materials, conductive carbon black, and polyvinylidene fluoride are used as effective ingredients.

[0090] In one preferred embodiment, the solid content of the electrode slurry is 35-45 wt %.

[0091] Based on the same inventive concept, the present invention also claims protection for a lithium battery negative electrode material, including a current collector, wherein an active material layer prepared from the electrode slurry is provided on the surface of the current collector.

[0092] In one preferred embodiment, the method for preparing the negative electrode material of the lithium battery is to uniformly apply the electrode slurry to the current collector and control the surface density of the active material to be 12-15 mg / cm 2 , and dried to obtain.

[0093] In one preferred embodiment, the current collector is a copper foil current collector.

[0094] In one preferred embodiment, the drying process is: pre-drying at 70-80°C for 1-2 hours, and then treating in a vacuum environment at 120-130°C for 10-14 hours.

[0095] In one preferred embodiment, the vacuum degree of the vacuum environment is ≤-0.08 MPa.

[0096] Based on the same inventive concept, the present invention also claims protection for a liquid battery, which includes the lithium battery negative electrode material.

[0097] In one preferred embodiment, in the liquid battery, the electrolyte is a mixed solution of ethylene carbonate and dimethyl carbonate containing 1-2M lithium hexafluorophosphate (LiPF6), with a volume ratio of ethylene carbonate / dimethyl carbonate of 1-2:1-2, and 2-3wt% fluoroethylene carbonate (FEC) and 1-2wt% vinylene carbonate (VC) are added as functional additives.

[0098] In one preferred embodiment, the liquid battery is LiNi 0.8 Co 0.1 Mn 0.1O2 (NCM811) is the positive electrode and polyethylene is the separator.

[0099] Based on the same inventive concept, the present invention also claims protection for a solid-state battery comprising the composite coating material.

[0100] Based on the same inventive concept, the present invention also claims protection for the application of the composite coating material in PVDF-HFP-based solid-state batteries.

[0101] In one preferred embodiment, in a PVDF-HFP-based solid-state battery, the electrolyte comprises 5-25 wt% PVDF-HFP, 5-25 wt% HFP, 5-20 wt% lithium lanthanum zirconium amide, 20-50 wt% lithium trifluoromethanesulfonyl imide, and 0.5-1.0 wt% of a crosslinker;

[0102] In PVDF-HFP, the mass ratio of PVDF to HFP is 80-90:15.

[0103] In a preferred embodiment, the cross-linking agent is selected from any one of peroxides, sulfur systems or silane coupling agents.

[0104] In one preferred embodiment, in a PVDF-HFP-based solid-state battery, a hexagonal boron nitride (h-BN) buffer layer is sprayed on the electrode interface.

[0105] In one preferred embodiment, the thickness of the hexagonal boron nitride (h-BN) buffer layer is 5±1 nm.

[0106] In one preferred embodiment, the density of the hexagonal boron nitride (h-BN) buffer layer is 0.5-1.2 mg / cm 2 .

[0107] In one preferred embodiment, the room temperature ionic conductivity of the PVDF-HFP based solid-state battery is ≥3.3×10 - 4 S / cm, 1C cycle 1000 times capacity retention rate ≥85%.

[0108] In one preferred embodiment, the preparation method of lithium lanthanum zirconium oxyamide includes: immersing hydroxylated lithium lanthanum zirconium oxy particles in 3-aminopropyltriethoxysilane ethanol solution, and refluxing at 80-90° C. for 4-6 hours to obtain lithium lanthanum zirconium oxyamide.

[0109] In one preferred embodiment, the concentration of the 3-aminopropyltriethoxysilane ethanol solution is 5-20 wt %.

[0110] In one preferred embodiment, the amino density of lithium lanthanum zirconium amide is 1.8-2.2 / nm2 .

[0111] In one preferred embodiment, the spraying process of the h-BN buffer layer adopts aerosol deposition technology with a carrier gas pressure of 0.2-0.5 MPa.

[0112] In a preferred embodiment, the peroxide cross-linking agent is dicumyl peroxide (DCP) or lauroyl peroxide (LPO).

[0113] In one preferred embodiment, the sulfur system is composed of sulfur (S8) and accelerator N-tert-butyl-2-benzothiazolesulfenamide (TBBS) in a mass ratio of 3-7:1.

[0114] In one preferred embodiment, the silane coupling agent is bis-(γ-triethoxysilylpropyl) tetrasulfide (Si69).

[0115] In one preferred embodiment, the electrolyte is formed into a film with a thickness of 20-100 μm by a tape casting method, and a gradient drying process is performed: pre-drying at 50° C. for 2 hours and then heating to 80° C. for 4 hours.

[0116] In one preferred embodiment, to improve interfacial compatibility, a hexagonal boron nitride (h-BN) nanosheet buffer layer with a thickness of 5±1 nm is introduced at the electrode / electrolyte interface, and its surface density is controlled to 0.5-1.2 mg / cm by aerosol spraying technology. 2 .

[0117] In one preferred embodiment, the final assembled full battery uses NCM811 as the positive electrode (area capacity 3.2 mAh / cm 2 ), SiC@polymer-LLZO as negative electrode (area capacity 4.0mAh / cm 2 ), the packaging is completed under 5-8MPa pressure conditions to achieve stable solid-state interface contact.

[0118] The present invention significantly improves the electrochemical performance and mechanical stability of silicon-carbon anode materials through composite coating structure design and interface coordinated regulation, and achieves dual adaptation of liquid and solid-state battery systems. Compared with existing technologies, the technical advantages of the present invention are specifically reflected in the following aspects:

[0119] 1. Multiple performance optimization of composite coating structure

[0120] The three-dimensional ion / electron dual continuous transport network constructed based on the gradient carbonization process makes the room temperature ionic conductivity of the composite coating material ≥2.3×10 -4S / cm, four orders of magnitude higher than traditional PAN-coated materials. Simultaneously, the buffering effect of the "macropore-mesopore-micropore" synergistic porous structure (macropores of 200-1200 nm, mesopores of 50-100 nm, and micropores of 5-15 nm) suppresses the silicon volume expansion rate from over 300% to 121% (Example 1). After 500 cycles at 1C, the capacity retention rate is 85.7% (Example 2). Using a low-temperature gradient carbonization process instead of traditional high-temperature sintering reduces energy consumption while avoiding grain coarsening and significantly improving batch consistency (Examples 1 and 3).

[0121] 2. Enhanced interface synergy and improved stability

[0122] The interfacial binding energy is increased to 9.2 J / m through the La-OC covalent bond formed by the surface hydroxyl groups of lithium lanthanum zirconium oxide and the polymer carbonization network. 2 (Example 2) significantly improved compared to the pure polymer-coated system (Comparative Example 1). In a PVDF-HFP-based solid-state system, the introduction of a 5±1 nm h-BN buffer layer improved room-temperature ionic conductivity. The assembled full battery exhibited enhanced interfacial impedance after 50 cycles and capacity retention after 500 cycles at 1C at 80°C. Specifically for solid-state battery systems, the fluorine-carbon bond interaction between the amino-modified LLZO filler and the PVDF-HFP matrix reduced interfacial impedance compared to the unmodified system.

[0123] 3. Cross-system application performance verification and industrial adaptation

[0124] The present invention exhibits an initial coulombic efficiency of 83.5% and a capacity retention rate of 88.8% after 500 cycles at 1C in liquid batteries, and is applicable to PVDF-HFP and PEO-based systems (Example 6 ionic conductivity 2.5×10 -4 S / cm). In solid-state battery applications, by adding 20-50% LiTFSI lithium salt to PVDF-HFP-based solid electrolytes, the room temperature ionic conductivity is increased to 3.39×10 -4 S / cm, and maintained dendrite-free performance after 500 cycles at 1C when paired with an NCM811 cathode at 80°C (Example 4). In industrial implementation, the material preparation process is highly compatible with existing production lines, and combined with solvent recovery, unit costs are reduced, demonstrating feasibility for large-scale production.

[0125] 4. Improved safety performance

[0126] The thermal runaway starting temperature of the composite coating material is increased to 200°C (Example 1), which is significantly improved compared to 128°C of the conventional silicon-carbon negative electrode (Comparative Example 1), proving that the safety performance of the composite coating material is greatly improved. BRIEF DESCRIPTION OF THE DRAWINGS

[0127] Figure 1 This is the SEM morphology of LLZO nanoparticles after gradient ball milling.

[0128] Figure 2 This is the SEM image of the cross section of the composite coating.

[0129] Figure 3 This is the XPS O1s spectrum of LLZO surface hydroxylation.

[0130] Figure 4 This is the temperature-time curve of the gradient carbonization process.

[0131] Figure 5 This is the 1C cycle performance curve of liquid battery electrodes.

[0132] Figure 6 This is the DSC test curve of the composite coating material of Example 1.

[0133] Figure 7 The relationship between the molecular weight of different porogens and the macroscopic pore size (logarithmic coordinates); in the figure, PEG fitting: D = 2.90M n 0.73 , PVP fitting: D = 5.33M n 0.51 .

[0134] Figure 8 This is the DSC test curve of the composite coating material of Comparative Example 1. DETAILED DESCRIPTION

[0135] The present invention is not limited to the following specific embodiments. Based on the disclosure of the present invention, a person skilled in the art may adopt a variety of other specific embodiments to implement the present invention. Any simple changes or modifications made to the design structure and concept of the present invention fall within the scope of protection of the present invention. It should be noted that the embodiments and features of the embodiments of the present invention may be combined with each other unless they conflict.

[0136] Example 1: Preparation of composite coating materials and verification of liquid battery performance

[0137] 10g Li 6.4 La3Zr 1.4 Ta 0.6 O 12 The powder was mixed with ethanol and subjected to gradient ball milling using zirconium oxide balls with diameters of 3 mm and 5 mm, and a volume ratio of 1:2. The speed of coarse grinding was 300 rpm for the first 2 hours and then switched to 150 rpm for fine grinding for the next 4 hours. The SEM morphology of LLZO nanoparticles after gradient ball milling is shown in Figure 2. Figure 1 The obtained powder was vacuum dried at 80℃ for 12 hours, then immersed in 5wt% nitric acid solution and ultrasonically treated at 80℃ for 2 hours, washed and dried to obtain pretreated LLZO. The X-ray photoelectron spectroscopy (XPS) analysis was performed on it, and the results were as follows: Figure 3 As shown, the surface hydroxyl density is 3.2 / nm 2 .

[0138] 8g PAN was dissolved in 92g N,N-dimethylformamide (DMF), stirred at a constant temperature of 60℃ for 6 hours to form a homogeneous solution, and 80g silicon-carbon core material (silicon mass accounts for 30wt%, D50 is 10μm, G4S-C800, Shanghai Shanshan) was added, followed by 2g pre-treated LLZO and 0.8g PEG-400. After ball milling and dispersion for 2 hours, gradient carbonization treatment was performed: under argon protection, the temperature was raised to 250℃ at 2℃ / min for pre-oxidation for 2 hours, and then raised to 800℃ at 5℃ / min for carbonization for 2 hours, with an argon flow rate of 4L / min. Finally, the temperature was kept at 650℃ for 15 minutes and slowly cooled to room temperature to obtain a composite coating material (SiC@PAN-LLZO composite powder). The temperature-time curve of the gradient carbonization process is shown in Figure 2. Figure 4 As shown. The cross section of the composite coating material was SEM-graphed, and the results were as follows: Figure 2 As shown in the figure, the macropore size is about 400 nm, the mesopore size is 50-100 nm, and the micropore size is 2-5 nm.

[0139] During the preparation process, specific porogens achieve precise construction of macroscopic pores; the gradient ball milling process (dynamic adjustment from 300 to 150 rpm) controls the D50 of the pretreated LLZO to 80 nm and optimizes the 50-100 nm mesopore distribution; through critical phase transition temperature control at 650°C, the graphitization degree of the carbon layer and the micropore structure (5-15 nm) are regulated to achieve a balance between ion / electron dual channels and mechanical stability.

[0140] In the composite coating material, macropores account for 50-60%. Three-dimensional reconstruction statistics were obtained using focused ion beam scanning electron microscopy (FIB-SEM), and pore volume was extracted using threshold segmentation. Mesopores account for 30-40%. Based on the BJH model of nitrogen adsorption-desorption isotherms, LLZO nanoparticles (D50 = 80-90 nm) were accumulated to form a sphere-like close-packed pore structure (porosity ≈26%). Electron tomography (ET) was used to verify the spatial uniformity of the distribution. Micropores account for 5-10%. Small-angle X-ray scattering (SAXS) combined with density functional theory (DFT) model fitting was used to quantify the lattice-level porosity caused by carbonization shrinkage (micropore density saturates when the carbon layer has a graphitization degree of ≥80%). Macropores preferentially absorb the macroscopic strain of silicon core expansion; mesopores disperse shear stress, preventing crack propagation; and micropores stabilize the SEI film through surface adsorption. Macropores are mainly distributed on the outer surface of the coating layer; mesopores are located in the LLZO particle accumulation area; and micropores are concentrated in the carbonized polymer matrix. The pore distribution basically conforms to the distribution pattern of sparse outside and dense inside (macropore density increases outward). DSC testing of the composite coating material shows that the thermal runaway starting temperature of the SiC@PAN-LLZO composite powder material increases to 200℃ ( Figure 6 ).

[0141] The obtained SiC@PAN-LLZO composite powder was mixed with Super P and PVDF in a mass ratio of 7:2:1. The slurry was prepared by using NMP as solvent and a double planetary stirring process (revolution 30 rpm / rotation 1200 rpm) for 4 hours. After coating on the copper foil current collector, the slurry was dried in a gradient manner to form a surface density of 12-15 mg / cm 2 The electrode was prepared using a gradient drying process: drying at 80°C for 2 h, then at 120°C for 12 h (80°C / 2 h → 120°C / 12 h) (vacuum degree ≤ -0.08 MPa). A 1M LiPF6 electrolyte in EC / DMC (volume ratio 1:1) was used, with 2 wt% FEC and 1 wt% VC added. NCM811 was used as the positive electrode (areal capacity 3.2 mAh / cm 2 ), and assembled the battery with a 9 μm thick PE separator.

[0142] Electrochemical tests show that the electrode has an initial coulombic efficiency of 83.5% at a rate of 0.1C and a capacity retention of 88.8% after 500 cycles at 1C (see Figure 5 The volume expansion rate was determined to be 118% by in-situ XRD, and the interface impedance was stable at 18.3Ω·cm 2 .

[0143] Test methods and data verification:

[0144] 1. Ionic conductivity: Calculated using the Chenhua CH760e electrochemical workstation using the formula σ = L / (R·A). The test frequency range is 0.1 Hz to 1 MHz.

[0145] 2. Volume expansion rate: In situ testing was performed using a Rigaku Smartlab SE X-ray diffractometer to refine and analyze changes in unit cell parameters.

[0146] 3. Interface binding energy: Nanomechanical mapping was performed using a Bruker edge atomic force microscope, and quantitative analysis was performed using force-displacement curves.

[0147] 4. Cycling performance: Using Xinwei BTS-5V10mA test system, the voltage window is set to 0.01-1.5V (vs. Li + / Li).

[0148] 5. Calculation of LLZO surface hydroxyl density:

[0149] The O 1s peak was analyzed using a Thermo Scientific K-Alpha X-ray photoelectron spectrometer with monochromatized Al Kα radiation (1486.6 eV) as the excitation source. The peak area attributable to hydroxyl (OH) (binding energy 532.1 ± 0.2 eV) was determined by peak fitting and calculated according to the formula:

[0150]

[0151] in, is the area ratio of the hydroxyl peak, is the total sensitivity factor, is the specific surface area of ​​LLZO (m² / g), d is the XPS detection depth, is Avogadro's constant.

[0152] 6. Calculation of LLZO surface amino group density:

[0153] Fitting the amino group (- ) characteristic peak (binding energy 399.6±0.2 eV), according to the formula:

[0154] ;

[0155] in is the ratio of the amino peak area, the coefficient 2 is the XPS quantitative calibration factor (based on standard sample calibration), and the other parameters are defined in the same way as the hydroxyl density calculation. + Sputter clean the surface for 30 seconds (1 keV, 1 μA / cm 2 ).

[0156] Example 2

[0157] PVDF-HFP based solid-state battery applications

[0158] Preparation of amino-modified LLZO nanoparticles: The hydroxylated LLZO nanoparticles in Example 1 were amino-modified by immersing the surface hydroxylated LLZO in a 3 wt% 3-aminopropyltriethoxysilane (APTES) ethanol solution and refluxing at 80°C for 6 hours to obtain a surface amino group density of 2.0 / nm. 2 LLZO nanoparticles.

[0159] Preparation of the solution system: PVDF-HFP containing 5-25 wt% hexafluoropropylene (HFP) (PVDF:HFP mass ratio of 85:15) was selected as the matrix material. 20-50 wt% lithium bis(trifluoromethanesulfonyl)imide (LiTFSI) was added as the lithium salt and dissolved in a polar organic solvent (DMF / acetone (7:3, v / v)) to form a solution. 3 g of amino-modified LLZO nanoparticles was then added to the solution, along with 0.12 g of a crosslinker, dicumyl peroxide (DCP), to form a composite electrolyte solution.

[0160] Battery assembly: The composite electrolyte solution was cast into a 20-100 μm thick film using a tape casting method. A gradient drying process was then implemented: pre-drying at 50°C for 2 hours followed by heating to 80°C for 4 hours. To improve interfacial compatibility, a 5±1 nm thick hexagonal boron nitride (h-BN) nanosheet buffer layer was introduced at the electrode / electrolyte interface. Its surface density was controlled to 1.0 mg / cm by aerosol spraying. 2 The test showed that the room temperature ionic conductivity of the electrolyte membrane reached 3.39×10 -4 S / cm.

[0161] The production of the positive and negative electrodes of the battery is the same as in Example 1. The initial interface impedance of the assembled full battery is 31.4Ω·cm 2 After 50 cycles, the interface impedance is 30.8Ω·cm 2 , the interface impedance is 37.9Ω·cm after 1000 cycles 2 , growth rate ≤ 5% ( Figure 7 The capacity retention rate after 1000 cycles at 2C is 85.2%, and the capacity retention rate after 500 cycles at 1C at 80°C is 84.7%, with no dendrite formation.

[0162] Example 3

[0163] Performance optimization of high LLZO content composites

[0164] The amount of LLZO added was adjusted to 3 g (PAN 7 g), and the rest of the process was the same as in Example 1. Performance tests showed that the ionic conductivity was increased to 3.1×10 -4 S / cm, the volume expansion rate further decreased to 105%, and the capacity retention rate was 85.7% after 500 cycles at 1C. Atomic force microscopy (AFM) mechanical testing showed that the interfacial bonding energy between LLZO and PAN carbonized network was enhanced to 9.2 J / m 2 , confirming the strengthening effect of La-OC covalent bond.

[0165] Example 4

[0166] Control of critical phase transition in gradient carbonization process

[0167] The carbonization procedure was changed to carbonization at 700° C. for 2 hours, followed by slow cooling to 550° C. The remaining processes were the same as in Example 1. After 500 1C cycles, the capacity retention rate was 82.3%.

[0168] Example 5

[0169] PEO-based composite coating materials

[0170] 8g PEO (M v = 600,000) was dissolved in THF / acetone (7:3, v / v), and 2 g Nanoparticles and 1.2g PEG-400, the rest of the process is the same as Example 1. The test shows that the ionic conductivity of the composite coating material is 2.3×10 -4 S / cm, and the volume expansion rate is 121%.

[0171] Example 6

[0172] Comparison of Al-doped LLZO

[0173] use Alternative The rest of the process is the same as in Example 1. The ionic conductivity is 2.1×10 -4 S / cm, and the capacity retention rate after 500 cycles is 86.5%.

[0174] Example 7

[0175] Optimization of PVP pore former

[0176] The porogen was replaced with 10 wt% polyvinylpyrrolidone (PVP K30), and the rest of the process was the same as in Example 4. SEM showed that the macropore size was 900 ± 75 nm. The capacity retention rate after 600 cycles at 1C was 90.2%. The relationship curve between the molecular weight of the PVP porogen and the macropore size is shown in Figure 7 .

[0177] Comparative Example 1

[0178] Pure PAN coating system without LLZO

[0179] The LLZO addition step was omitted, and the rest of the process was the same as in Example 1. The test results showed that the volume expansion rate increased to 276%, and the capacity retention rate was only 73.2% after 100 cycles at 1C, and it decayed rapidly in subsequent cycles ( Figure 5 ), which is significantly inferior to the solution of Example 1 of the present invention, confirming the core role of LLZO in ion transport and mechanical buffering. According to DSC testing, the thermal runaway temperature of the material is 128℃ ( Figure 8 ), the security of the certification material is also inferior to that of the solution in Example 1 of the present invention.

[0180] Comparative Example 2

[0181] The hydroxyl density of hydroxylated LLZO is less than 3 active sites per nm²

[0182] 10g The powder was subjected to gradient ball milling under the same conditions as in Example 1. The resulting powder was immersed in a 1 wt% nitric acid solution and ultrasonically treated at 80°C (frequency 40 kHz, power 300 W) for 30 minutes. XPS analysis showed that the surface hydroxyl density was only 1.5 / nm. 2 The remaining processes are the same as in Example 1.

[0183] Electrochemical tests showed that the battery's initial coulombic efficiency dropped to 72.3%, and the capacity retention rate after 500 cycles at 1C was only 58.6%. In-situ XRD measurements showed that the volume expansion rate of the silicon-carbon negative electrode material was 215%, and the initial interfacial impedance was 185Ω·cm 2 , after 500 cycles at 1C, it increased to 320Ω·cm 2 , much higher than Example 1.

[0184] Comparative Example 3

[0185] Fixed argon flow rate

[0186] During the carbonization process, the argon flow rate was fixed at 2 L / min, and the rest of the process was the same as in Example 1. SEM showed that the pore connectivity was poor and the ionic conductivity of the composite coating material was 1.1×10 -4 S / cm, and the capacity decays to 75% after 300 cycles.

[0187] Comparative Example 4

[0188] Single ball milling parameters

[0189] The ball milling was carried out at a single speed of 300 rpm for 5 hours (without gradient adjustment). The LLZO particle size D50 was 150 nm. The rest of the process was the same as in Example 1. SEM showed that the mesopore size was greater than 200 nm and the ionic conductivity was 1.8×10 -4 S / cm (22% lower than in Example 1), the volume expansion rate increased to 185%, and the active material after cycling was clearly pulverized and broken by SEM. The open porosity was less than 70% (mercury intrusion porosimetry). The battery's capacity retention after 500 1C cycles was only 75.3%, lower than in Example 1.

[0190] The initial value of the interface impedance is 85Ω·cm 2 , increased to 152Ω·cm after 500 cycles at 1C 2 .

[0191] Comparative Example 5

[0192] Lack of macropores

[0193] The porogen PEG-400 was omitted, and the remaining experimental steps were the same as in Example 1.

[0194] The pore structure lacks macropores and only has mesopores and micropores, and the rest is the same as in Example 1.

[0195] The results show that the volume expansion rate of the silicon-carbon negative electrode material is >200% as determined by in-situ XRD.

[0196] Comparative Example 6

[0197] The pre-oxidation step was omitted in the gradient heat treatment, and the temperature was directly increased to 800° C. at a rate of 10° C. / min for carbonization for 3 hours. The remaining processes were the same as in Example 1.

[0198] The pore structure lacks micropores and only has mesopores and macropores, and the rest is the same as in Example 1.

[0199] The results show that after 500 cycles of 1C, the capacity attenuation rate of the battery increased by 30% compared with Example 1.

[0200] Comparative Example 7

[0201] Composite coating materials containing only macroscopic pores (200-1000nm)

[0202] The ball milling process was adjusted to a single speed of 400 rpm for 5 hours, and the LLZO particles had a D50 of 120 nm. The pre-oxidation step was omitted from the gradient heat treatment, and the temperature was directly increased to 800°C at a rate of 10°C / min for carbonization for 3 hours. The remaining processes were the same as in Example 1.

[0203] SEM showed that the pores consisted of only 200-1000 nm macropores (mesopores / micropores accounted for <5%). The volume expansion rate reached 253% (in situ XRD), and the capacity retention rate after 500 cycles at 1C was only 62.1%. The initial interfacial impedance value was 85 Ω·cm 2 , increased to 210Ω·cm after cycling 2 (AFM shows crack growth.) The tests showed that a single macroscopic pore was unable to disperse multi-scale stresses, confirming the necessity of multi-level pore coordination.

[0204] Comparative Example 8

[0205] Macropore diameter 100nm (<200nm)

[0206] PEG-200 (molecular weight 200) was selected as the porogen, and the addition amount was 15 wt %. The rest was the same as in Example 1.

[0207] The average macropore diameter was tested to be 100±15nm (mercury intrusion porosimetry), with a normal ratio of mesopores to micropores. The volume expansion rate rebounded to 182% (silicon particle expansion was not adequately buffered), and the electrode pulverization rate was >30% after 300 cycles at 1C. The ionic conductivity was 1.7×10 -4 S / cm (26% lower than Example 1). The tested structures showed that insufficient macropore size resulted in decreased buffering efficiency.

[0208] Comparative Example 9

[0209] Macropore diameter 1200nm

[0210] PEG-4000 (molecular weight 4000) was used as a porogen in an amount of 20 wt %. The carbonization temperature was raised to 850° C., and the rest was the same as in Example 1.

[0211] The macropore size was tested to be 1200 ± 50 nm (SEM analysis), with the mesopores and micropores squeezed. The electrode sheet breakage rate during coating was +20% (due to loose structure caused by excessive pores). After 200 cycles at 1C, the capacity retention rate was 71.3%, and the interfacial impedance was >150 Ω·cm. 2 The test results show that excessively large pores weaken the mechanical strength, proving that 1000 nm is the upper limit of the size.

[0212] Comparative Example 10

[0213] Do not use the gradient cooling mode, and directly cool down to room temperature at a rate of 5-7℃ / min.

[0214] After carbonization, the 650°C holding period was cancelled and the product was directly cooled to room temperature. The remaining process parameters were the same as those in Example 1. The performance comparison between Comparative Example 10 and Example 1 is shown in the table below.

[0215] Table 1 Performance comparison between comparative example 10 and embodiment 1

[0216]

[0217] The lack of a 650°C holding period reduces the degree of aromatization of the carbon chain, resulting in a lower proportion of micropores. This reduction in graphitization also leads to a decrease in the material's electronic conductivity. This also reduces the graphite layer's ability to inhibit thermal decomposition.

[0218] Comparative Example 11

[0219] The gradient heat treatment was adjusted to: heating at 10°C / min to 1000°C for carbonization for 2 hours, and eliminating the 650°C holding step.

[0220] The graphitization degree was tested to be >90% (XRD analysis), and the micropores collapsed (SAXS showed a 60% decrease in micropore density). The ionic conductivity dropped to 1.1×10 -4 S / cm (the carbon layer is too dense to prevent Li + The volume expansion rate rose back to 168% (porous structure failure). The test results show that excessively high temperatures destroy the porous structure.

[0221] Comparative Example 12

[0222] The carbonization temperature was set to 500° C., and the rest was the same as in Example 1.

[0223] The polymer was tested and found to be incompletely carbonized (Raman ID / IG = 1.8, compared to 0.9 in Example 1). Electronic conductivity < 10 - 3 S / cm (four-probe method), 1C capacity is only 820 mAh / g (1350 mAh / g in Example 1). Volume expansion is >200% (the coating layer does not form an effective buffer). Test results indicate that low-temperature carbonization leads to conductive network defects.

[0224] Comparative Example 13

[0225] Gradient-free heat treatment

[0226] The gradient temperature increase program was cancelled, and the temperature was directly increased from room temperature to 800° C. at a rate of 10° C. / min for carbonization for 2 h.

[0227] The pore connectivity was poor (mercury intrusion porosity <70%), and the initial interface impedance was 83 Ω·cm. 2 After 500 cycles at 1C, the capacity retention rate was 75.6% (compared to 88.8% in Example 1). The thermal runaway temperature dropped to 175°C (DSC test, due to heat accumulation caused by uneven pore distribution). Test results show that gradient heat treatment is crucial for oriented pore arrangement, and direct temperature increase leads to performance degradation.

[0228] Comparative Example 14

[0229] Referring to the prior art CN 112467116 A, a graphite core (silicon mass percentage 0%) and an unhydroxylated LLZO coating were used. Tests showed: volume expansion >300% (silicon unbuffered), capacity retention of only 65% ​​after 100 cycles at 1C, and interface impedance >200 Ω·cm. 2 , confirming that the graphite core cannot adapt to the high expansion characteristics of silicon-based materials.

[0230] Comparative Example 15

[0231] Referring to the prior art CN 111244410 B, a multi-layer coating (carbon layer + HF insulation layer + SEI layer) was used, and the gradient dispersion process was omitted. The results showed that the ionic conductivity of the SEI layer was only 1.2×10 -6 S / cm, the capacity retention rate after 500 cycles is 73.2%, and the interface delamination causes the electrode powdering rate to be >40%.

[0232] Comparative Example 16

[0233] Referring to the prior art CN 115241526 A, lithium borate (2wt%) was used to replace amino LLZO. Tests showed that the expansion suppression efficiency was only 38%, and the electron mobility dropped to 10 -4 S / cm, and the capacity decayed to 68% after 300 cycles at 1C, proving that lithium borate cannot adapt to silicon-based stress.

[0234] It should be noted that the above embodiments are merely examples for the purpose of clearly illustrating the present invention and are not intended to limit the embodiments of the present invention. A person skilled in the art would be able to make other variations or modifications based on the above description. It is not possible to enumerate all embodiments here. Any obvious variations or modifications arising from the technical solution of the present invention remain within the scope of protection of the present invention.

Claims

1. A composite coating material, characterized in that: include: The core is a core-shell structure; the shell structure of the core is made of carbon, and the core structure of the core is made of silicon-based material; as well as A shell coated on the outer surface of the core, the shell comprising a porous structure, and containing a polymer and hydroxylated lithium lanthanum zirconium oxide particles; Among them, the structural formula of lithium lanthanum zirconium oxide particles is Li 7-x La3Zr 2-y M y O 12 , M is at least one of the doping elements Ta, Al, and Nb that replaces Zr in equal amounts, 0≤x≤1, 0≤y≤0.6; The pore structure includes macropores with a pore diameter of 200-1000 nm, mesopores with a pore diameter of 50-100 nm, and micropores with a pore diameter of 2-5 nm; During the preparation of the composite coating material, a gradient heat treatment is performed; during the gradient heat treatment, the polymer forms a carbonized network; and the hydroxylated lithium lanthanum zirconium oxide particles and the polymer carbonized network form La-OC covalent bonds.

2. The composite coating material according to claim 1, characterized in that: In the pore structure, macropores account for 50-60%, mesopores account for 30-40%, and micropores account for 5-10%.

3. The composite coating material according to claim 1, characterized in that The mass of the core structure accounts for 15-50wt% of the core; the silicon-based material includes elemental silicon, porous silicon, nano silicon, and the general formula SiO x , a combination of one or more silicon-oxygen materials with 0<x<2; the thickness of the shell structure is 50-200 nm, and the crystallinity is ≥75%; the material of the shell structure is graphitized carbon.

4. The composite coating material according to claim 1, characterized in that: The particle size D50 of the core is 5-12 μm, and the specific surface area is 1-3 m 2 / g, and the tap density is 0.8-1.2 g / cm 3 The porosity of the core is ≤5%; and / or the surface of the hydroxylated lithium lanthanum zirconium oxide particles has a hydroxyl density greater than 3 / nm 2 active site.

5. The composite coating material according to claim 1, characterized in that: The preparation method of hydroxylated lithium lanthanum zirconyl oxide particles comprises the following steps: ball milling and drying the lithium lanthanum zirconyl oxide particles, and then adding them to an acid solution for ultrasonic treatment to obtain hydroxylated lithium lanthanum zirconyl oxide particles; wherein the ball milling is performed using a gradient dispersion process, and the gradient dispersion process is as follows: first, ball milling at a rotation speed of 300-400 rpm for 1-3 hours, and then ball milling at a rotation speed of 100-150 rpm for 3-5 hours.

6. The method for preparing the composite coating material according to any one of claims 1 to 5, characterized in that: The following steps are involved: S1. Mixing a polymer and hydroxylated lithium lanthanum zirconium oxide particles in a mass ratio of 8-5:2-5; adding a porogen after mixing evenly, and sonicating to obtain a mixture; S2, mixing the mixture and the core uniformly to obtain a pre-product; S3, performing gradient heat treatment on the pre-product to obtain a composite coating material; The gradient heat treatment process is as follows: under argon protection, heating to 200-250°C at a rate of 2-3°C / min for pre-oxidation for 1-2 hours, then heating to 700-800°C at a rate of 5-10°C / min for carbonization for 2-3 hours, then cooling to 600-650°C at a rate of 5-7°C / min for insulation for 15-20 minutes, and then slowly cooling to room temperature.

7. The preparation method according to claim 6, characterized in that In step S1, the amount of the porogen added is 0.1-15wt% of the polymer mass; the porogen is one or both of polyethylene glycol and polyvinyl pyrrolidone; in step S2, the mass ratio of the core to the polymer in the mixture is 10:0.5-10:3; in step S3, the slow cooling stage is carried out at a rate of 10-15°C / min to cool to room temperature.

8. An electrode slurry, characterized in that: The electrode slurry includes the composite coating material according to any one of claims 1 to 5.

9. A lithium battery negative electrode material, characterized in that The present invention comprises a current collector, wherein an active material layer prepared from the electrode slurry according to claim 8 is provided on the surface of the current collector.

10. A battery, characterized in that: The battery is a liquid battery, and the liquid battery includes the lithium battery negative electrode material according to claim 9; or, the battery is a solid-state battery, and the solid-state battery includes the composite coating material according to any one of claims 1 to 5.

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