A lithium-rich manganese-based material and its application in sulfide all-solid-state batteries

Abstract

The application discloses a lithium-rich manganese-based material, which has an inner core and a coating layer, wherein the inner core is Li 1+ x TM 1‑x Ru y O2, TM is selected from one or more of Mn, Ni and Co, 0

Description

Technical Field

[0001] This invention relates to the field of materials technology, specifically to a lithium-rich manganese-based material for use in sulfide all-solid-state batteries. Background Technology

[0002] Rechargeable batteries are a key technology for clean energy and electric vehicle applications. However, traditional lithium-ion batteries (LIBs) use flammable liquid electrolytes and graphite anodes (which reduce energy density compared to lithium metal anodes). As an alternative, all-solid-state lithium-ion batteries (ASSLBs) using solid electrolytes (SEs) are gaining increasing attention due to their superior safety, wide electrochemical stability window, and the potential to achieve high energy densities. Therefore, all-solid-state lithium-ion batteries using solid electrolytes are strong candidates for next-generation energy storage devices.

[0003] Meanwhile, some traditional cathode materials, such as LiCoO2 and nickel-rich oxide cathode materials, have been successfully applied to sulfide-based all-solid-state lithium batteries. However, due to the inherent capacity limitations of the materials themselves, the capacity of most cathode materials is limited to 200 mAh g⁻¹. -1 Unlike traditional cathode materials whose capacity derives solely from the redox reactions of transition metal (TM) cations, lithium-rich manganese-based layered oxide (LLO) cathode materials derive their ultra-high capacity from multi-electron redox reactions, including both cation and anion redox reactions (the cation redox reaction is typically Ni). 2+ / Ni 4+ Co 3+ / Co 4+ Small amount of Mn 3+ / Mn 4+ Anionic redox reactions are usually O 2- / O - / O2), thus providing a larger capacity (>250mAh g) than conventional cation redox. -1 This can break through the capacity bottleneck that stems solely from traditional TM-centered redox activities.

[0004] However, there are still technical challenges in applying lithium-rich manganese-based layered oxide cathode materials to sulfide all-solid-state lithium batteries. First, there are inherent problems of lithium-rich manganese-based layered oxide cathode materials: First, the anion redox process of lithium-rich manganese-based layered oxide cathode materials easily leads to irreversible oxygen precipitation, especially in the surface region, which reduces the initial Coulombic efficiency (ICE) and hinders the application of LLO. Particularly when working under a high-voltage environment (≥4.8V), irreversible oxygen precipitation leads to the generation of too many oxygen vacancies in the structure of lithium-rich manganese-based layered oxide, thereby resulting in structural collapse of the lithium-rich manganese-based layered oxide; Second, excessive spinel phase on the surface of lithium-rich manganese-based oxide reduces the capacity of the LLO cathode. Secondly, there are technical problems that need to be solved when applying lithium-rich manganese-based layered oxide cathode materials to sulfide all-solid-state lithium batteries: First, the lithium-rich manganese-based layered oxide cathode material will undergo severe interfacial side reactions with the sulfide solid electrolyte; Second, an interfacial space charge layer (SCL) is formed between the lithium-rich manganese-based layered oxide cathode material and the sulfide solid electrolyte due to the difference in electrochemical potential, which affects the transport of lithium ions in the battery; Third, due to hindered lithium ion transport, the activation of the Li₂MnO₃ phase during the first charging of the battery is severely limited, thereby impeding the specific capacity of lithium-rich manganese-based layered oxide in sulfide all-solid-state lithium batteries. Fourth, the irreversible oxygen precipitation of the lithium-rich manganese-based layered oxide cathode material is more significant in sulfide solid electrolyte batteries. Summary of the Invention

[0005] In view of the problems in the prior art, the present invention discloses a lithium-rich manganese-based material. When the lithium-rich manganese-based material of the present invention is used in sulfide solid electrolyte batteries, it can deliver a higher specific capacity, and also enables the sulfide solid electrolyte battery to have better cycling performance.

[0006] The present invention is achieved through the following technical solutions:

[0007] The present invention provides a lithium-rich manganese-based material, said lithium-rich manganese-based material (LLO-Ru@S) has a core and a coating layer, wherein the core is Li 1+x TM 1-x Ru y O₂, TM is selected from one or more of Mn, Ni and Co, 0<x<0.3, 0<y<0.02, and the coating layer is a sulfur-containing coating layer.

[0008] The above-described design of this invention involves Ru doping into lithium-rich manganese-based materials. Ru forms a strong covalent Ru-O bond with O, enhancing the stability of the oxygen lattice and reducing the generation of excessive oxygen vacancies, thus preventing structural collapse. The sulfur-containing coating layer not only isolates the sulfide solid electrolyte from surface side reactions with the lithium-rich manganese-based material, but also further improves the stability of the oxygen lattice in the material. Based on this, Ru doping can reduce the spinel phase on the surface of lithium-rich manganese-based materials, while the sulfur-containing coating layer can further promote the increase of the rock salt phase on the surface of lithium-rich manganese-based materials. With the combined effect of Ru and S in the sulfur-containing coating layer, the diffusion ability of lithium ions in lithium-rich manganese-based materials can be enhanced, and the diffusion ability of lithium ions in the Li2MnO3 phase of lithium-rich manganese-based materials can be enhanced, thereby stimulating the activation of the Li2MnO3 phase, which is beneficial to the specific capacity of lithium-rich manganese-based materials, and also enhances the lithium ion diffusion ability on the surface of lithium-rich manganese-based materials, effectively alleviating the interfacial space charge layer.

[0009] The lithium-rich manganese-based material LLO-Ru@S in this invention. Here, "-" indicates that Ru element is doped into the lithium-rich manganese-based crystalline oxide; "@" indicates that the material following "@" coats the material preceding "@".

[0010] As a further option, the raw material for the sulfur-containing coating layer includes CS2 (carbon disulfide).

[0011] As a further option, the Ru element is derived from one or more of Ru₂O (ruthenium dioxide), ruthenium chloride, and ruthenium acetylacetonate.

[0012] As a further option, the thickness of the sulfur-containing coating layer is 8nm-12nm.

[0013] As a further embodiment, the lithium-rich manganese-based material has a layered phase and a rock salt phase.

[0014] As a further improvement, the thickness of the rock salt phase in the lithium-rich manganese-based material is not less than 2.5 nm.

[0015] The present invention also provides a method for preparing the lithium-rich manganese-based material, the method comprising: weighing lithium source, TM source and Ru source according to elemental stoichiometry, grinding and mixing them evenly, and performing three-stage sintering to obtain a lithium-rich manganese-based material precursor (LLO-Ru); adding CS2 and performing heat treatment to obtain the lithium-rich manganese-based material of the present invention.

[0016] As a further option, the lithium source includes one or more of Li2CO3 (lithium carbonate), lithium manganese oxide, lithium iron phosphate, lithium acetate, and lithium nitrate.

[0017] As a further option, the TM source includes one or more of Co source, Mn source, and Ni source.

[0018] As a further option, the Co source includes one or more of cobalt sulfate, cobalt nitrate, cobalt chloride, nickel acetate, and manganese-rich ternary carbonates.

[0019] As a further option, the Mn source includes one or more of manganese sulfate, manganese nitrate, manganese chloride, manganese acetate, and manganese-rich ternary carbonates.

[0020] As a further option, the Ni source includes one or more of nickel sulfate, nickel nitrate, nickel chloride, nickel acetate, and manganese-rich ternary carbonates.

[0021] As a further option, the TM source can be selected from the same type, including manganese-rich ternary carbonates.

[0022] As a further embodiment, the manganese-rich ternary carbonate includes (Ni x Co y Mn 1-x-y One or more of CO3, wherein x>0, y>0, and x+y<1.

[0023] As a further solution, the (Ni) x Co y Mn 1-x-y CO3 includes Ni 0.16 Co 0.16 Mn 0.68 CO3.

[0024] As a further option, the Ru source includes one or more of Ru2O, ruthenium chloride, and ruthenium acetylacetonate.

[0025] As a further step, the three sintered sections need to be cooled and broken down after sintering. Those skilled in the art can choose the cooling method according to the actual situation, such as natural cooling.

[0026] As a further step, the heat treatment needs to be followed by cooling. Those skilled in the art can choose the cooling method according to the actual situation, such as natural cooling.

[0027] As a further step, the lithium-rich manganese-based precursor needs to be crushed before adding CS2. This helps to expose the lithium-rich manganese-based precursor, thereby ensuring that the coating layer can fully coat the surface of the precursor.

[0028] As a further embodiment, the lithium source is added in excess by 5%. This excess lithium source is necessary to compensate for the reduction in lithium due to the volatilization of lithium carbonate during sintering. In this invention, the 5% excess lithium source means that, based on the stoichiometric calculation of the required mass of lithium source, it is added in excess by 5%.

[0029] As a further embodiment, the grinding and mixing time is 0.8h-1.2h. Those skilled in the art can adjust the grinding and mixing time according to actual conditions, as long as sufficient homogeneity between the materials is ensured.

[0030] As a further embodiment, the three-stage sintering specifically includes a first stage with a sintering temperature of 300℃-400℃ and a sintering time of 4h-6h, a second stage with a sintering temperature of 680℃-720℃ and a sintering time of 14h-16h, and a third stage with a sintering temperature of 840℃-880℃ and a sintering time of 4h-6h.

[0031] As a further embodiment, the heating rate of the three-stage sintering is 2.5℃ / min-3.5℃ / min.

[0032] As a further embodiment, the temperature of the heat treatment is 150℃-250℃, and the heat treatment time is 0.8h-1.2h.

[0033] As a further option, the heat treatment conditions also include a nitrogen atmosphere.

[0034] The present invention also provides a positive electrode or electrochemical device having the aforementioned lithium-rich manganese-based material.

[0035] As a further embodiment, the electrochemical device can be used in end-consumer products, including but not limited to mobile phones, laptops, pen input computers, mobile computers, e-book players, portable phones, portable fax machines, portable copiers, and portable printers.

[0036] As a further embodiment, the electrochemical device can be used in electrical equipment, including large and small electrical equipment. Small electrical equipment includes consumer products, wearable electronic devices, or portable electronic devices; large electrical equipment includes transportation equipment. Transportation equipment includes, but is not limited to, vehicles such as automobiles, motorcycles, electric bicycles, buses, subways, high-speed trains, airplanes, and ships. Wearable electronic devices or portable electronic devices include, but are not limited to, headphones, video recorders, LCD TVs, portable cleaners, portable CD players, mini CDs, transceivers, electronic notebooks, calculators, memory cards, portable recorders, radios, backup power supplies, drones, motors, lighting fixtures, toys, game consoles, clocks, power tools, flashlights, cameras, large household batteries, and lithium-ion capacitors.

[0037] As a further embodiment, the electrochemical device includes a battery.

[0038] As a further option, the battery includes an all-solid-state ion battery.

[0039] As a further embodiment, the all-solid-state ion battery includes a sulfide all-solid-state electrolyte battery. The lithium-rich manganese-based material of the present invention is particularly suitable for use in sulfide all-solid-state electrolyte batteries, where it can improve the battery's discharge specific capacity and cycle performance.

[0040] As a further embodiment, the sulfide all-solid-state electrolyte battery also includes a negative electrode and a sulfide solid-state electrolyte.

[0041] As a further option, the negative electrode is Li-In.

[0042] As a further option, the negative electrode active material in the negative electrode is Li4Ti5O. 12 .

[0043] As a further option, the sulfide solid electrolyte includes one or more of the following: LISICON type sulfide solid electrolyte, LGPS type sulfide solid electrolyte, and sulfosilver germanite type solid electrolyte.

[0044] The LISICON-type sulfide solid electrolytes include: (100-x)Li₂S-xP₂S₅, (100-x)Li₂S-xSiS₂, and Li₄-yGe. 1-y P y One or more of S4, where 0 < x < 100, 0 < y < 1;

[0045] The general chemical formula of the LGPS type sulfide solid electrolyte is Li 11-z M 2-z P 1+z S12 , where 0 < z < 2, and M includes one of the elements Ge, Si, and Sn;

[0046] The chemical formula of the sulfosilver germanite-type solid electrolyte is Li6PS5X, where X includes one of the elements Cl, Br, and I.

[0047] As a further preferred option, the sulfide solid electrolyte is Li6PS5Cl.

[0048] The features and beneficial effects of this invention are as follows:

[0049] (1) The lithium-rich manganese-based material obtained by the present invention has a more stable structure. Applying the lithium-rich manganese-based material of the present invention to sulfide solid electrolyte batteries can improve the cycle performance and discharge specific capacity of sulfide solid electrolyte batteries.

[0050] (2) The lithium-rich manganese-based material of the present invention, with the combined use of Ru doping and sulfur-containing coating, can not only stabilize lattice oxygen and improve the structural stability of the lithium-rich manganese-based material, but also construct ion diffusion channels in the lithium-rich manganese-based material. On this basis, the lithium-rich manganese-based material of the present invention will not undergo side reactions with sulfide solid electrolytes, and can also effectively alleviate SCL reactions, thereby improving the electrical performance of the lithium-rich manganese-based material. Attached Figure Description

[0051] To more clearly illustrate the technical solutions in the embodiments of the present invention, the drawings used in the description of the embodiments or the prior art will be briefly introduced below. Obviously, the drawings described below are only some embodiments of the present invention. For those skilled in the art, other drawings can be obtained based on these drawings without creative effort.

[0052] Figure 1 This is a schematic diagram illustrating the strategy of the lithium-rich manganese-based material of the present invention.

[0053] Figure 2 This is a schematic diagram of the element distribution in the embodiments and comparative examples of the present invention, wherein... Figure 2 a is the LLO's STEM-EDS element mapping; Figure 2 b is the XRD pattern of LLO, LLO-Ru, LLO@S and LLO-Ru@S; Figure 2 c is a magnified view of the XRD pattern of LLO-Ru.

[0054] Figure 3 These are STEM images and corresponding atomic structures of phase boundaries observed in the materials of the embodiments and comparative examples of the present invention, wherein... Figure 3 a represents the comparative LLO; Figure 3 b is a proportional LLO-Ru; Figure 3c is the embodiment LLO-Ru@S.

[0055] Figure 4 The embodiments and comparative examples of this invention demonstrate the electrochemical performance when applied in sulfide solid electrolyte batteries, wherein... Figure 4 a represents the charge-discharge curves of the LLO and LLO-Ru@S samples; Figure 4 b is the cyclic voltammetry curve of the embodiments and comparative examples of the present invention at 0.05 mV / s; Figure 4 c is the cyclic voltammetry curve of the LLO-Ru@S sample for the first four cycles; Figure 4 d represents the cycling performance of the embodiments and comparative examples of the present invention at 0.1C; Figure 4 e represents the long-cycle performance of the LLO-Ru@S sample at 0.1C; Figure 4 f is the charge-discharge curve of LLO and LLO-Ru@S samples at 55℃; Figure 4 g represents the dQ / dV curves for LLO and LLO-Ru@S samples; Figure 4 h represents the cycling performance of the embodiments and comparative examples of the present invention at 0.5C; Figure 4 i represents the rate capability of the embodiments and comparative examples of the present invention; Figure 4 j represents the cycling performance of the LLO-Ru@S sample at a discharge rate of 1C.

[0056] Figure 5 The impedance of the embodiments and comparative examples of the present invention when applied in sulfide solid electrolyte batteries is shown, wherein Figure 5 a is the impedance spectrum of the LLO / LPSCl / Li-In battery at different cycle numbers; Figure 5 b is the impedance spectrum of the LLO-Ru / LPSCl / Li-In battery at different cycle numbers; Figure 5 c is the impedance spectrum of the LLO@S / LPSCl / Li-In battery at different cycle numbers; Figure 5 d is the impedance spectrum of the LLO-Ru@S / LPSCl / Li-In battery at different cycle numbers; Figure 5 e is the DRT curve calculated by the environmental impact system in LLO / LPSCl / Li-In batteries; Figure 5 f is the DRT curve calculated by the environmental impact system in the LLO-Ru / LPSCl / Li-In battery; Figure 5 g is the DRT curve calculated by the environmental impact system in the LLO@S / LPSCl / Li-In battery; Figure 5 h is the DRT curve calculated by the environmental impact system in LLO-Ru@S / LPSCl / Li-In battery; Figure 5i represents the impedance spectrum of LLO / LPSCl / Li-In cell, LLO-Ru / LPSCl / Li-In cell, LLO@S / LPSCl / Li-In cell and LLO-Ru@S / LPSCl / Li-In cell after one cycle; Figure 5 j is the impedance spectrum of LLO / LPSCl / Li-In cell, LLO-Ru / LPSCl / Li-In cell, LLO@S / LPSCl / Li-In cell and LLO-Ru@S / LPSCl / Li-In cell after 5 cycles; Figure 5 k is the impedance spectrum of the LLO / LPSCl / Li-In cell, LLO-Ru / LPSCl / Li-In cell, LLO@S / LPSCl / Li-In cell, and LLO-Ru@S / LPSCl / Li-In cell after 10 cycles. Figure 5 l represents the impedance spectrum of LLO / LPSCl / Li-In cell, LLO-Ru / LPSCl / Li-In cell, LLO@S / LPSCl / Li-In cell, and LLO-Ru@S / LPSCl / Li-In cell after 30 cycles. Figure 5 m is the DRT curve calculated based on LLO / LPSCl / Li-In battery, LLO-Ru / LPSCl / Li-In battery, LLO@S / LPSCl / Li-In battery and LLO-Ru@S / LPSCl / Li-In battery after one cycle; Figure 5 n is the DRT curve calculated based on LLO / LPSCl / Li-In battery, LLO-Ru / LPSCl / Li-In battery, LLO@S / LPSCl / Li-In battery and LLO-Ru@S / LPSCl / Li-In battery after 5 cycles; Figure 5 o is the DRT curve calculated based on LLO / LPSCl / Li-In battery, LLO-Ru / LPSCl / Li-In battery, LLO@S / LPSCl / Li-In battery and LLO-Ru@S / LPSCl / Li-In battery after 10 cycles; Figure 5 p is the DRT curve calculated based on LLO / LPSCl / Li-In, LLO-Ru / LPSCl / Li-In, LLO@S / LPSCl / Li-In, and LLO-Ru@S / LPSCl / Li-In batteries after 30 cycles.

[0057] Figure 6 The examples and comparative examples of this invention are CV curves applied to sulfide solid electrolyte batteries at different scan rates, wherein... Figure 6 a is the CV curve of LLO at different scan rates; Figure 6 b is the CV curve of LLO-Ru at different scan rates; Figure 6 c is the CV curve of LLO@S at different scan rates; Figure 6 d represents the CV curves of LLO-Ru@S at different scan rates; Figure 6 e is the peak current (ip) in the CV curve and the square root of the scan rate (V). 1 / 2 The relationship between ) Figure 6 f is the conductivity of the embodiments and comparative examples of the present invention.

[0058] Figure 7 The XPS spectra of the positive electrode in a sulfide solid electrolyte battery under different states are shown in the embodiments and comparative examples of the present invention. Figure 7 a is the S2p XPS spectrum of the LLO electrode and the LLO-Ru@S electrode under different conditions (no cycling, after 70 cycles); Figure 7 b is the O1s XPS spectrum of the LLO electrode and LLO-Ru@S electrode under different conditions (no cycling, after 70 cycles).

[0059] Figure 8 Here is a TEM-EDS elemental mapping of the comparative LLO@S, where, Figure 8 a is the SEM image of LLO@S. Figure 8 b is the distribution diagram of the Mn element. Figure 8 c is the distribution diagram of Ni element. Figure 8 d is the distribution map of Co element. Figure 8 e is the distribution diagram of element S.

[0060] Figure 9 Here is a comparative LLO-Ru TEM-EDS elemental mapping, where, Figure 9 a is the SEM image of LLO-Ru. Figure 9 b is the distribution diagram of the Mn element. Figure 9 c is the distribution map of Ni element. Figure 9 d is the distribution map of Co element. Figure 9 e is a distribution diagram of the Ru element.

[0061] Figure 10 Here is the TEM-EDS element mapping diagram of the LLO-Ru@S example, where, Figure 10 a is the SEM image of LLO-Ru@S. Figure 10 b is the distribution diagram of the Mn element. Figure 10 c is the distribution map of Ni element. Figure 10 d is the distribution map of Co element. Figure 10 e is a distribution diagram of the Ru element. Figure 10 f is the distribution diagram of element S.

[0062] Figure 11 Here are high-resolution transmission electron microscopy (HRTEM) images and Fast Fourier Transform (FFT) modes of the LLO@S, in which Figure 11 a represents the thickness of the sulfur-containing coating layer, which is 9.92 nm. Figure 11 b represents the interplanar spacing of 0.47 nm, corresponding to the (003) crystal plane. Figure 11 c is the diffraction spot obtained after Fourier transform.

[0063] Figure 12 The electrochemical performance of the LLO sample at room temperature is given by [the relevant data]. Figure 12 a is the charge / discharge curve; Figure 12 b is the cyclic voltammetry curve; Figure 12 c represents the cycling performance at 0.1C; Figure 12 d represents the rate capability at 0.1C, 0.3C, 0.5C, 1C, 2C, and 5C, respectively. Detailed Implementation

[0064] To facilitate understanding of the preparation method of the lithium-rich manganese-based material of the present invention, a more comprehensive description of the preparation method of the lithium-rich manganese-based material of the present invention will be given below, and embodiments of the present invention will be provided, but this does not limit the scope of the present invention.

[0065] In this invention, Ni from manganese-rich ternary carbonates is used. 0.16 Co 0.16 Mn 0.68 Taking CO3 as an example, further research was conducted to prepare lithium-rich manganese-based materials for use in batteries, which improved the battery's electrical performance.

[0066] Example 1: Using Li2CO3 and Ni 0.16 Co 0.16 Mn 0.68 CO3 and Ru2O are Ru sources; Ni is weighed according to stoichiometry. 0.16 Co 0.16 Mn 0.68CO3, Ru2O, and Li2CO3, with Li2CO3 added in excess of 5%, were mixed using an agate mortar or ball mill for 1 hour, followed by calcination at 350°C for 5 hours, 700°C for 15 hours, and 860°C for 5 hours in air (heating rate = 5°C / min). After natural cooling to room temperature, LLO-Ru core material was obtained. LLO-Ru was then treated by grinding the LLO-Ru core material in a mortar for 30 minutes. LLO-Ru was then placed in an alumina crucible and transferred to a tube furnace. Next, 80 mL of liquid CS2 (99.9%, Aladdin) was placed in a bottle. Due to the low boiling point of CS2, it readily converted from the liquid phase to the gas phase and was blown into the tube furnace by a N2 gas stream. LLO-Ru was then synthesized into powder by heating at 200°C for 1 hour in an atmosphere of N2 and CS2 mixed gas. After natural cooling to room temperature, LLO-Ru@S lithium-rich manganese-based material was obtained.

[0067] Comparative Example 1: Using Li2CO3 and Ni 0.16 Co 0.16 Mn 0.68 CO3, Ni was weighed according to stoichiometry. 0.16 Co 0.16 Mn 0.68 CO3 and Li2CO3, with Li2CO3 added in excess of 5%, were mixed using an agate mortar or ball mill for 1 hour, followed by calcination at 350°C for 5 hours, 700°C for 15 hours, and 860°C for 5 hours in air atmosphere (heating rate = 5°C / min). LLO material was obtained after natural cooling to room temperature.

[0068] Comparative Example 2: Using Li2CO3 and Ni 0.16 Co 0.16 Mn 0.68 CO3 and Ru2O are Ru sources; Ni is weighed according to stoichiometry. 0.16 Co 0.16 Mn 0.68 CO3, Ru2O, and Li2CO3, with Li2CO3 added in excess of 5%, were mixed using an agate mortar or ball mill for 1 hour. The mixture was then calcined in air at 350°C for 5 hours, 700°C for 15 hours, and 860°C for 5 hours (heating rate = 5°C / min). After natural cooling to room temperature, LLO-Ru material was obtained.

[0069] Comparative Example 3: Using Li2CO3 and Ni 0.16 Co 0.16 Mn 0.68 CO3, Ni was weighed according to stoichiometry. 0.16 Co 0.16 Mn 0.68CO3 and Li2CO3, with Li2CO3 added in excess of 5%, were mixed using an agate mortar or ball mill for 1 hour, followed by calcination at 350°C for 5 hours, 700°C for 15 hours, and 860°C for 5 hours in air (heating rate = 5°C / min). LLO material was obtained after natural cooling to room temperature. The LLO material was then treated by grinding in a mortar for 30 minutes. The LLO was then placed in an alumina crucible and transferred to a tube furnace. Next, 80 mL of liquid CS2 (99.9%, Aladdin) was placed in a bottle. Due to the low boiling point of CS2, it readily converted from the liquid phase to the gas phase and was blown into the tube furnace by a N2 gas stream. The LLO was then heated to 200°C for 1 hour in an atmosphere of N2 and CS2 mixed gas to obtain powder, which was then naturally cooled to room temperature to obtain LLO@S material.

[0070] The preparation method of the sulfide solid electrolyte (Li6PS5Cl solid electrolyte) in this invention is as follows: LiCl (lithium chloride) (Sigma-Aldrich; >99%), P2S5 (phosphorus pentasulfide) (MACKLIN; 99%), and Li2S (lithium sulfide) (Sigma-Aldrich; >99%) are weighed according to stoichiometry, then mixed. The mixture is ball-milled at 300 rpm for 48 hours in a 250 mL zirconia jar using a mixture of zirconia balls with diameters of 10 mm, 8 mm, and 5 mm, with a ball-to-powder ratio of 40:1. The recovered powder is ground and pressed into granules using a mortar and pestle. The granules are sealed in a glass tube and annealed at 550°C for 10 hours. The resulting solid electrolyte has an ionic conductivity of approximately 3.5 mS / cm at room temperature. -1 The sulfide solid electrolyte in this invention is also commercially available.

[0071] We also obtained lithium-rich manganese-based cathode materials from the examples and comparative examples for use in sulfide solid-state electrolyte batteries. The negative electrode of the battery used Li-In, with Li6PS5Cl (LPSCl) as the sulfide solid electrolyte, and a laboratory-scale all-solid-state battery with a diameter of 10 mm was prepared. All assembly processes were performed in a dry, argon-filled glove box. For the cathode, sulfide solid electrolyte, lithium-rich manganese-based cathode material, and conductive agent (VGCF) were selected and mixed in agate slurry at a mass ratio of 50:45:5 (lithium-rich manganese-based cathode material:LPSCl:VGCF) for 1 hour to obtain a cathode composite powder. The mass of the sulfide solid electrolyte was 80 mg, and VGCF represents carbon fiber.

[0072] First, a dense sulfide solid electrolyte layer was prepared by pressurizing 80 mg of sulfide solid electrolyte (Li6PS5Cl) at 240 MPa. Then, 5 mg of positive electrode composite powder was uniformly coated onto one side of the sulfide solid electrolyte layer, and a pressure of 360 MPa was applied. Next, an In foil with a diameter of 10 nm and a thickness of 50 μm was pressed onto the other side, followed by a thin lithium foil with a diameter of 8 nm and a thickness of 30 μm. Finally, the cell was fixed in a stainless steel casing under a pressure of approximately 20 MPa.

[0073] Preparation method of LLO / Li liquid half-cell: A slurry of lithium-rich manganese-based cathode material (90 wt%), carbon black (5 wt%), and PVDF (polyvinylidene fluoride) binder (5 wt%) was coated onto Al foil. After uniform stirring, the stable slurry was coated onto conductive aluminum foil and dried at 120°C for 8 hours. The dried electrode was cut into 12 mm diameter discs as working electrodes. Swagelok-type half-cells were assembled in an argon-filled glove box, with a lithium metal disc as the counter electrode and a glass microfiber filter using the liquid electrolyte as a separator.

[0074] This invention will also obtain lithium-rich manganese-based materials for testing:

[0075] (1) X-ray diffraction (XRD) patterns were obtained on a Bruker AXSD8 Advance, with Cu Kα radiation in the range of 10° ≤ 2θ ≤ 80°.

[0076] Atomic structures were characterized by a TEM (JEOL, JEM-F200) operating at 200 kV, and HADDF-STEM was characterized by a Cs-corrected STEM (FEI Titan Cubed Themis G2300) operating at 300 kV.

[0077] X-ray photoelectron spectroscopy (XPS) measurements were performed using a monochromatic Al-Kα (1486.6 eV) X-ray source on a Thermo ESCALAB 250 system to investigate the relative abundance and chemical state of elements. The XPS data were fitted using the asymmetric Gaussian-Lorentz sum function in XPSPEAK software and the atomic relative abundance calculations in Multipak software.

[0078] (2) The galvanic charge-discharge test and cycle performance of the assembled all-solid-state LIB were evaluated using a standard battery testing instrument (LAND CT-2001A, Wuhan Langbo Test Equipment Co., Ltd.) at room temperature and 55°C. All tests were performed at a current rate of 0.1C (1C = 200mAh) between 1.4 and 4.2V (compared to Li-In). The selected test program included stepwise increases in the C-type rate, increasing by five cycles at a time.

[0079] (3) Discharge / charge tests of the liquid half-cell were performed on a Land BT2000 battery test system (Wuhan, China) at a current density of 0.1C, with a voltage window of 2.0 to 4.8V (relative to Li / Li). + Cyclic voltammetry (CV) tests were performed on an electrochemical workstation (CHI660E, Shanghai, China) at a rate of 0.1 mV / s. -1 The frequency sweep is performed.

[0080] (4) Solid-state battery testing was performed on a Land BT2000 battery testing system (Wuhan, China) at a current density of 0.1C, with a voltage window of 1.4 to 4.2V (vs. Li-In). Cyclic voltammetry (CV) testing was conducted on an electrochemical workstation (CHI660E, Shanghai, China) at a current density of 0.03mVs. -1 0.05mVs -1 0.07mVs -1 0.1mVs -1 0.2mVs -1 The frequency sweep is performed.

[0081] Verification Result Analysis:

[0082] Table 1

[0083]

[0084] The method of this invention successfully prepared lithium-rich manganese-based materials. We applied these materials, along with comparatively prepared lithium-rich manganese-based materials, to sulfide solid-state electrolyte batteries. We found that the lithium-rich manganese-based material of this invention significantly improved the battery's discharge specific capacity and exhibited good cycle performance. We believe this is partly due to the doping of Ru in this invention. Ru forms Ru-O bonds with O in the LLO (Lithium-ion Loop) electrolyte, which have strong covalent bond characteristics. These Ru-O bonds stabilize the lattice oxygen in the LLO, greatly improving the structural stability of the lithium-rich manganese-based material and preventing structural collapse caused by excessive oxygen vacancies, thus affecting lithium-ion transport and deintercalation. Furthermore, the shell structure of this invention is a sulfur-containing coating layer. This sulfur-containing coating layer not only isolates the interfacial side reactions between the sulfide solid electrolyte and the lithium-rich manganese-based material but also further enhances the stability of the lattice oxygen in the structure. Furthermore, the sulfur coating layer can interact with Ru. First, Ru doping reduces the spinel phase on the LLO surface, which is beneficial for obtaining an LLO with only the rock salt phase. The sulfur-containing coating layer promotes the increase of the rock salt phase, thereby improving the electrical performance of lithium-rich manganese-based materials. Second, the interaction between S and Ru in the sulfur-containing coating layer can enhance the lithium-ion diffusion capacity of lithium-rich manganese-based materials. On the one hand, it can activate the Li2MnO3 phase in lithium-rich manganese-based materials, which is beneficial for the capacity utilization of lithium-rich manganese-based materials. On the other hand, it can effectively alleviate the interfacial space charge layer, thereby improving the electrical performance of lithium-rich manganese-based materials. Therefore, we further studied and compared to verify that the lithium-rich manganese-based materials obtained by the method of the present invention can be applied to sulfide solid electrolyte batteries and have good electrical performance.

[0085] To address the challenges of applying lithium-rich manganese-based layered oxide cathode materials in sulfide-based all-solid-state lithium batteries, we propose a dual modification strategy that addresses both internal and external factors. This represents a crucial step in the application of lithium-rich manganese-based layered oxide cathode materials in sulfide-based all-solid-state lithium batteries. For example... Figure 1 As shown, a trace amount of ruthenium is first introduced to obtain the doped cathode material Li. 1.2 Ni 0.13 Co 0.13 Mn 0.54 Ru 0.01 O2(LLO-Ru), trace doping of Ru enhances the Li + The diffusion ability of Ru in the Li₂MnO₃ phase then stimulates the activation of the Li₂MnO₃ phase. Simultaneously, Ru doping in the cathode material Li… 1.2 Ni 0.13 Co 0.13 Mn 0.54 Ru 0.01In O2, a strong covalent Ru-O bond is formed with O, enhancing the stability of the lattice oxygen and greatly avoiding the formation of interfacial side reaction products during the electrochemical reaction between the sulfide solid electrolyte and LLO. Based on this, a time- and cost-effective sulfidation strategy is used to deposit an ultrathin (8nm-12nm) and uniform sulfur-containing coating layer on the surface of the LLO-Ru cathode material. This not only constructs a stable interface but also significantly promotes the Li-O exchange at the LLO-Ru and sulfide solid electrolyte interface. + The transport properties of LLO-Ru@S can effectively mitigate interfacial side reactions between LLO and sulfide solid electrolytes. Simultaneously, the low electronic conductivity of LLO-Ru@S alleviates the SCL effect. The results show that the modified LLO-Ru@S exhibits excellent electrochemical performance, including satisfactory rate capability (140.6 mAh g⁻¹ at 0.5C). -1 This results in an ultra-long lifetime of up to 2022 cycles at 4.2V (vs. Li-In) in sulfide all-solid-state lithium batteries (capacity retention >70% at 1C). Importantly, a thorough and detailed analysis was conducted, revealing the fundamental reasons for the improved interfacial stability and lithium diffusion kinetics of lithium-rich materials in sulfide all-solid-state lithium batteries. Through targeted improvements, interfacial degradation caused by oxygen release under high voltage was suppressed at its source, opening up new avenues for the application of lithium-rich manganese-based materials in high-energy-density sulfide solid-state electrolyte batteries.

[0086] Based on this, we further compare the characteristics of different modified lithium-rich manganese-based materials obtained from the embodiments and comparative examples of the present invention. The preparation processes of Comparative Example 1 (lithium-rich layered oxide (LLO), Comparative Example 2 (ruthenium-doped layered oxide (LLO-Ru), Comparative Example 3 (sulfide layered oxide (LLO@S)) and Example 1 (sulfide layered oxide (LLO-Ru@S)) are as follows. Figure 2 As shown in Figure a, the scanning transmission electron microscopy-energy dispersive X-ray spectroscopy (STEM-EDS) spectrum of LLO shows that Ni, Co, and Mn are uniformly distributed in the LLO particles in the embodiments and comparative examples of this invention. The energy spectrum of LLO@S (as shown in Figure a) Figure 8 As shown in the figure, the distribution of S overlaps with the distribution of Ni, Co, and Mn elements, confirming that S is uniformly distributed on the surface of LLO particles. Figures 9-10 It was also shown that trace amounts of Ru and S were uniformly distributed in LLO-Ru particles and LLO-Ru@S particles, respectively. The structural features of all embodiments and comparative examples are as follows: Figure 2 As shown in b, the X-ray diffraction (XRD) results of the four samples show that the α-NaFeO2 hexagonal layered phase structure with R-3m space group in the embodiments and comparative examples of the present invention can be classified as a Li2MnO3-like phase (space group C2 / m) structure at the superlattice peak (red line) at 20°-30°. Figure 2The results in b show that the LLO structures in all four samples contain two phases. Furthermore, no additional impurity peaks were observed, indicating that neither the sulfur-containing coating nor Ru doping altered the crystal structure of LLO. To gain a deeper understanding of the structure and crystal information of LLO-Ru materials, the XRD spectra were refined using Rietveld refinement with a goodness-of-fit parameter χ² (2.04). Figure 2 c). Table 1 summarizes the detailed structural parameters after refinement. The c / a ratio of the LLO-Ru sample is 4.999, which is greater than the critical lattice constant of 4.899. This indicates that after Ru doping, the cations in the LLO-Ru lattice are still highly ordered, which also verifies that Ru doping has not changed the internal crystal structure of LLO.

[0087] The effects of sulfur-containing coatings and Ru doping on the surface morphology and structure of four samples were carefully investigated at the atomic scale using transmission electron microscopy (TEM) and high-angle annular dark-field scanning electron microscopy (HAADF-STEM). Figure 3 In step a, by identifying the boundaries in the LLO using atomic resolution STEM imaging, we found that the sample surface contains a small amount of spinel and rock salt phases of approximately 1 nm to 1.6 nm. The effect of sulfidation on surface structure modification was investigated using transmission electron microscopy. Figure 11 A sulfur-containing coating layer approximately 10 nm thick was observed on the surface of the LLO particles, the structure of which differs significantly from the bulk crystal region. The interplanar spacing of the LLO@S sample was measured to be approximately 0.49 nm, consistent with the (003) crystal plane of LLO. These findings were compared with XRD (… Figure 2 b) and TEM-EDS Figure 8 The combined results show that the LLO with a sulfur-containing coating was successfully constructed without affecting the bulk structure of the LLO cathode material. The sulfur-containing coating creates a stable interface, preventing direct contact between the lithium-rich layered oxide material (LLO) and the sulfide solid electrolyte, thereby suppressing interfacial side reactions. Figure 3 b is an atomic image of LLO-Ru, showing a disordered layered phase characteristic at the boundary (approximately 1 nm) between the rock salt phase and the layered phase. This result indicates that the introduction of Ru affects the spinel and rock salt phases on the LLO surface, and Ru doping can reduce the spinel phase on the LLO surface. An atomic image of the LLO-Ru@S example is shown below. Figure 3 As shown in c, due to surface modification and bulk doping, more rock salt phase with a thickness of 3 nm was formed on the LLO surface. This increased rock salt phase is beneficial for activating more Li₂MnO₃ phase, thereby improving the specific capacity of lithium-rich manganese-based matrix oxide and sulfide all-solid-state batteries. It is evident that the combined effect of Ru and the sulfur-containing coating layer not only reduces the spinel phase but also promotes the increase of the rock salt phase, which may greatly improve the electrical performance of lithium-rich manganese-based matrix oxide and sulfide all-solid-state batteries.

[0088] We used the embodiments and comparative examples of this invention as positive electrode materials in sulfide solid-state electrolyte batteries, wherein the negative electrode of the sulfide solid-state electrolyte battery is a Li-In negative electrode. In sulfide all-solid-state batteries ( Figure 4 a) At room temperature, the initial charge-discharge specific capacities of the LLO-Ru@S cathode material and the LLO cathode material are 113.8 mAh g, respectively. -1 and 54.8mAhg -1 The initial discharge specific capacity of the cathode material obtained by this invention is significantly higher than that of the comparative LLO cathode material. Based on this, we... Figure 4 In comparison (LLO), the charging plateau for lithium ion extraction from LiTMO2 is significantly shortened, while oxygen ion activation in the Li2MnO3 phase exhibits almost no charging plateau in the sulfide solid electrolyte battery. This indicates that lithium ion transport is severely impeded during the initial charging of the comparative example (LLO), limiting the activation of Li2MnO3. This is likely due to the high impedance caused by severe side reactions at the LLO / LPSCl interface, which hinders lithium ion transport. In comparison, the charging platforms of the LiTMO2 and Li2MnO3 phases corresponding to the LLO-Ru@S cathode are longer, and the overpotential of the LLO-Ru@S sulfide all-solid-state lithium battery is lower than that of LLO. This indicates that our work effectively enhances the lithium-ion transport dynamics. We believe this is due to two main reasons: First, Ru doping forms a strong covalent bond Ru-O with O atoms. The participation of this strong covalent bond enhances the oxygen lattice stability of LLO and suppresses the interfacial side reactions caused by oxygen release in the LLO sample, thus constructing a stable interface that is beneficial for ion migration. Second, the sulfide coating layer greatly improves the interfacial compatibility between the LLO cathode material and the sulfide solid electrolyte. CV curves at different scan rates confirm the higher Li-phase efficiency of LLO-Ru@S. + Diffusion coefficient (see) Figure 6 d). Figure 4 b shows a comparison of the CV curves of the four materials in the first cycle. The LLO-Ru@S sample in Example 1 exhibits stronger oxidation and reduction peaks, which can be attributed to better kinetic processes at the LLO-Ru@S / LPSCl interface. Notably, the oxidation peak of Li2MnO3 in the LLO-Ru@S sample is significantly stronger than that of the other three materials, further confirming the higher activation level of the Li2MnO3 phase in this invention. The CV curves of the LLO-Ru@S sample in the first four cycles of the sulfide solid electrolyte battery are shown below. Figure 4As shown in Figure c, two main redox peaks appear in the voltage regions of approximately 3.4 V (vs. Li-In) and 4.0 V (vs. Li-In). The initial oxidation peak at 3.4 V (vs. Li-In) corresponds to the oxidation reaction of TMs (transition metal cations), while the oxidation peak at 4 V (vs. Li-In) can be attributed to the activation process of Li₂MnO₃. After the first cycle, the oxidation peak at 4.0 V (vs. Li-In) almost disappears. Furthermore, no additional redox peaks for Ru were observed, possibly due to trace Ru doping. The cycling performance of the four materials was further evaluated at room temperature with a discharge rate of 0.1 C. Figure 4 d) Compared with pure LLO material, the initial specific discharge capacity of LLO-Ru@S in Example 1 is significantly improved. Figure 4 e represents the long-cycle performance of the LLO-Ru@S cathode at room temperature with a discharge rate of 0.1C. It can be seen that after 650 cycles, the capacity of the LLO-Ru@S sample showed almost no decay, indicating that the LLO-Ru@S and Li6PS5Cl exhibit good interfacial stability during electrochemical processes. We further investigated the ion dynamics performance of the embodiments and comparative examples at high temperatures, testing them at 55°C. Figure 4 As shown in f, the charge / discharge capacity of the comparative LLO at 55℃ is 209.8 mAh g. -1 and 119.9mAh g -1 As can be seen from the charging platform, the LiTMO2 phase is fully activated, indicating that high temperature is beneficial for activating the Li2MnO3 phase. However, the activation of the Li2MnO3 phase is still limited, due to the poor conductivity of Li2MnO3 and the incompatibility of the interface structure between LLO and the sulfide solid electrolyte. In stark contrast, the charging capacity of the LLO-Ru@S electrode in this embodiment of the invention reaches as high as 301.1 mAh g⁻¹. -1 The charging platform indicates that the low capacity of LLO in sulfide solid electrolyte batteries is mainly due to the limited activation of the Li2MnO3 phase. However, due to the dual modification strategy of Ru doping and sulfur-containing coating in this invention, good interfacial compatibility and stability promote bulk and interfacial ion transport, enabling the discharge capacity of LLO-Ru@S to reach 204 mAh g⁻¹. -1 The lithium-rich manganese-based material of this invention, when used in sulfide solid electrolyte batteries at high temperatures, exhibits electrical performance comparable to that of LLO in liquid electrolyte batteries. Figure 12 As shown.

[0089] exist Figure 4 The dQ / dV (differential capacity overvoltage) curve of g and Figure 4The initial discharge performance of f is in good agreement, which further indicates that more Li2MnO3 phases are activated in the LLO-Ru@S sample. Figure 4 h represents the cycling performance of the four materials at 0.5C. The initial specific discharge capacities of LLO, LLO-Ru, LLO@S, and LLO-Ru@S in the sulfide solid electrolyte are 45 mAh g. -1 76.7mAh g -1 94.1mAh g -1 and 140.6mAh g -1 This further demonstrates that Ru doping and sulfidation strategies can effectively improve the electrochemical performance of LLO in sulfide solid electrolytes. Figure 4 The results show the rate performance of four samples at different current densities. The results indicate that within the current density range of 0.1C-10C, the sulfidation strategy can establish a stable interface layer, reduce SCL, effectively suppress interfacial side reactions, and promote ion transport at the LLO / LPSCl interface, thus achieving a trend of LLO-Ru@S>LLO@S>LLO-ru>LLO. However, sulfidation alone cannot suppress oxygen release from the LLO material during cycling, while the covalent Ru-O bonds formed by Ru substitution can stabilize lattice oxygen, fundamentally preventing further oxygen-containing interface decomposition. Therefore, through dual modification of doping and surface, a good ion transport channel is constructed, which can accelerate lithium-ion diffusion, promote reaction kinetics, and thus help improve the rate performance. Furthermore, the sulfide solid-state electrolyte battery containing LLO-Ru@S (… Figure 4 j) It provides an ultra-long lifetime of 2022 cycles with a capacity retention of >70% at 1C, which further demonstrates the advantages of the combination of LLO-Ru@S and sulfide solid electrolyte.

[0090] Assessing interfacial stability and lithium-ion transport kinetics is crucial for explaining changes in electrochemical behavior in sulfide solid electrolyte batteries, which can be revealed through electrochemical impedance spectroscopy (EIS) measurements combined with relaxation time distribution (DRT) analysis. Figure 5 Impedance spectral variations of sulfide solid electrolyte batteries with four different cathode materials at 55°C are presented. EIS analysis was performed using a Zennium Pro electrochemical workstation within a frequency range of 0.01 Hz to 8 MHz, with an amplitude of 5 mV. The semicircle represents the charge transfer resistance (Rct), which primarily originates from the interfacial resistance between the cathode material and the sulfide solid electrolyte. Prior to testing, the batteries were charged to 4.2 V (vs. Li-In) through various cycles. Figure 5 a- Figure 5 Figure d shows the Nyquist plots of a sulfide solid-state electrolyte battery using four different cathode materials after 1, 5, 10, and 30 cycles. From Figure 5As can be seen, the Rct of the comparative LLO increases sharply from the 1st cycle to the 30th cycle (from 1600Ω to 3400Ω). This is due to severe interfacial chemical / electrochemical side reactions occurring between the LLO cathode and the LPSCl electrolyte during cycling. In stark contrast, the Rct value of the LLO-Ru@S sulfide solid electrolyte battery changes relatively little (e.g., ...). Figure 5 (As shown in d). It is noteworthy that after 30 cycles, the Rct of LLO-Ru@S remains lower than the initial Rct of the comparative LLO. This indicates that the LLO-Ru@S cathode material possesses stable and rapid interfacial diffusion kinetics during electrochemical processes. To accurately distinguish the evolution of Rct, the DRT technique was used to decouple the resistance evolution of the cathode material during lithium-ion insertion / deintercalation. The DRT fitting results are shown in d. Figure 5 As shown in e-5h, three main regions can be distinguished during cycling, labeled R1, R2, and R3. The R1 and R2 peaks represent the charge transfer resistance at the cathode material / LPSCl interface. Since LLO consists of two phases, it is reasonable for the two peaks observed in the DRT analysis to correspond to R1 and R2. The R3 peak is attributed to the solid-phase diffusion of lithium ions within the cathode material. Figure 5 As shown in Figure e, the charge transfer resistance at the interfaces (R1 and R2) dominates during the first cycle, while after 30 cycles, the value of R3 increases sharply, indicating that the high interfacial resistance between LLO and the sulfide solid electrolyte severely hinders the solid diffusion of ions in the LLO material. Figure 5 As can be seen from f, the values ​​of R1, R2, and R3 are basically equal after a large amount of Ru doping. This confirms that Ru doping not only improves the diffusion ability of Li in the Li2MnO3 composition, but also, due to the strong covalent bond between Ru and O, inhibits the decomposition of sulfide solid electrolytes related to oxygen release, which is beneficial to interface stability. Figure 5 As shown in g, the R1 and R2 values ​​of LLO@S are less than R3, indicating that the sulfidation strategy can construct a stable interface between LLO and LPSCl, effectively suppressing the large interfacial resistance. It is worth noting that existing technology reports indicate the presence of a significant amount of Li2MnO3 phase in LLO. This abundant Li2MnO3 phase leads to poor lithium-ion transport kinetics in sulfide solid-state electrolyte batteries. [1] The R1 and R2 values ​​of LLO-Ru@S after 30 cycles in this invention are... Figure 5 h) R1 and R2 values ​​lower than LLO-@S ( Figure 5 As can be seen from g), Ru doping and sulfidation strategies can enhance the lithium-ion transport dynamics of lithium-rich manganese-based materials. To more intuitively understand the effects of doping and surface modification on alleviating interfacial impedance, we compared the Nyquist curves of four materials after different cycle numbers, as shown in the figure. Figure 5 i- Figure 5 As shown in Figure l. The corresponding DRT fitting curve is shown in Figure l. Figure 5 m- Figure 5 As shown in p. Figure 5 Taking m as an example, it can be seen that LLO-Ru@S has the smallest R1, R2, and R3 values. It is worth noting that the R1, R2, and R3 values ​​of LLO-Ru are all lower than those of LLO, which further confirms that Ru doping can stabilize lattice oxygen, thereby eliminating interfacial degradation caused by oxygen release. Furthermore, compared with LLO and LLO-Ru, the R1 and R2 values ​​of LLO@S are significantly lower. This is because the sulfur-containing coating layer with low electronic conductivity suppresses SCL; at the same time, the sulfur-containing coating layer can also reduce interfacial side reactions caused by direct contact between LLO and LPSCl, thereby greatly improving the Li at the LLO / LPSCl interface inside the sulfide solid electrolyte battery. + Transportation. In short, Figure 5 The data shows that the LLO exhibits significant internal resistance during cycling. However, doping and sulfidation strategies can suppress interfacial side reactions, thereby helping to reduce impedance. These data confirm the benefits of interfacial stability in suppressing interfacial resistance.

[0091] Figure 6 a- Figure 6 d shows the CV curves of the examples and comparative examples at different scan rates. Figure 6 e describes the peak current (i p ) and the square root of the scan rate (V 1 / 2 The relationship between lithium-ion diffusion rate (D) and other parameters. Li+ The answer is obtained from the Randles-Sevcik equation.

[0092]

[0093] In the formula, n is the number of electrons in the redox reaction; A is the geometric area of ​​the electrode (cm²). 2 C0 is the molar concentration of lithium ions (mol / cm³). -3 V is the potential scan rate (Vs). -1 For the two samples in this paper, we can assume that n, A, and C0 are equal. Li+ This represents the diffusion coefficient of lithium ions in the layered structure, which is related to the scan rate v and i. p related. Figure 6 f shows i p For V 1 / 2The linear fitting results show that the diffusion rate of LLO-Ru@S is significantly higher than that of LLO. The low diffusion coefficient of LLO is related to oxygen release and interfacial degradation under high voltage, while doping and sulfidation create a stable interface, improving the migration rate of Li diffusion. The high interfacial resistance between the oxide cathode and the sulfide electrolyte is caused by the space charge layer formed at the interface. Due to the chemical voltage difference, the oxide cathode has a higher resistance to Li. + The attraction of is stronger than that of sulfide electrolytes, resulting in a small amount of Li + They migrate from the electrolyte surface to the oxide cathode side, forming a lithium-deficient space charge layer. For example... Figure 6 As shown in f, the LLO@S cathode material exhibits the lowest electronic conductivity, indicating that the sulfur-containing coating layer can significantly suppress the space charge layer in sulfide solid electrolyte batteries to a certain extent, thereby greatly improving electrochemical performance. The extremely low interfacial resistance also confirms this (see f). Figure 5 i). The difference is that the electronic conductivity of the LLO-Ru@S electrode is slightly higher than that of the LLO@S cathode, so its effect on mitigating the spatial interface layer is lower than that of the pure sulfide cathode, which also explains why its interface resistance is slightly higher than that of the LLO@S cathode.

[0094] Through surface-sensitive XPS measurements, we gained further detailed chemical insights into the side reactions occurring at the LLO / LPSCl interface under different states. It is important to note that, to maintain consistent testing conditions, the cathode material and LPSCl solid electrolyte (Li6PS5Cl solid electrolyte) were assembled into a complete battery, and then the cathode was disassembled for XPS testing. Figure 7 a represents the S2p XPS spectra of the LLO / LPSCl composite electrode before and after 70 cycles, and the S2p XPS spectrum of the LLO-Ru@S / LPSCl composite electrode after 70 cycles. Due to spin-orbit coupling, the s2p signal splits into two components with an area ratio of 2 / 1. Each chemical environment of S corresponds to a 2p3 / 2–2p1 / 2 heavy state. Figure 7 In Figure a, the S2p spectrum of the original LLO electrode shows a dominant doublet, with the S2ps3 / 2 component having a binding energy of 161.7 eV (blue component), which is the sulfur atom of LPSCl. The doublet at 160 eV can be attributed to Li2S, which comes from the reactants remaining during the LPSCl synthesis process. This means that the LLO cathode and LPSCl electrolyte are chemically stable before the start of electrochemical charge-discharge. After the 70th cycle, the two main peaks of LPSCl (blue component) weaken, and other components of the LLO / LPSCl system can be observed. The BE (doublet) at 162.1 eV (green component) in the S2p spectrum can be attributed to lithium polysulfides (Li2Sn), and the doublet at 164 eV (pink component) can be attributed to polysulfides (P2Sn). xLi2Sn and P2S x Byproducts from the sulfide solid electrolyte confirmed the interfacial instability between the LLO cathode material and the sulfide solid electrolyte. These byproducts hinder ion transport, increase interfacial resistance, and ultimately lead to capacity decay. In stark contrast, at the LLO-Ru@S / LPSCl cathode interface, the main peak intensity of LPSCl was strong, while the signal intensities of degradation products (~162.1 eV and ~164 eV) were weaker, and the relative content of byproducts was also reduced. This indicates that the sulfidation strategy is beneficial for LPSCl to maintain a reversible redox process and effectively suppresses interfacial reactions with the oxide cathode during cycling. The O1s spectra of the two cathodes in the XPS results are shown below. Figure 7 As shown in b. It should be noted that the experiment was conducted under vacuum, and the solid electrolyte did not contain oxygen. For the pristine LLO, lattice oxygen (O) 2- The corresponding O1s XPS peak is located at 529.3 eV. Furthermore, higher binding energy peaks are assigned to oxygen-containing materials deposited on the surface. After 70 cycles, the lattice oxygen O characteristic peak at 529.3 eV... 2- The voltage gradually decreases, and a new shoulder appears around 530.8 eV, which can be attributed to the oxidation of lattice oxygen (O). n- The formation of O2. 2- The lattice oxygen is extremely unstable and readily reacts with LPSCl electrolyte, which well explains the formation of interfacial side reaction products in the s 2p spectrum. However, higher TM-O bond covalentity may be beneficial for stabilizing O2. n- This indicates that at the LLO-Ru@S / LPSCl cathode interface, after 70 cycles, the peak intensity of 530.8 eV at the LLO-Ru@S / LPSCl cathode interface is weaker than that at the LLO / LPSCl cathode interface after 70 cycles, while O2... n- The relative content decreases. The oxygen stability of LLO affects its oxygen release behavior, and interfacial decomposition is directly related to interfacial oxygen loss. During battery cycling, lattice oxygen O2, which has high reactivity, is released. 2- The reaction with LPSCl electrolyte is violent, leading to the oxidation of LPSCl electrolyte. However, the highly stable Ru-O bonds generated by the abundant Ru doping fundamentally prevent further decomposition of the LPSCl electrolyte. Therefore, the highly stable interface LLO-Ru@S is derived from the co-modification of a sulfur-containing coating layer and Ru doping, achieving enhanced long-cycle stability under high voltage.

[0095] In summary, to meet the growing demands for high energy density and high safety, this invention utilizes the lithium-rich manganese-based cathode material and Li6PS5Cl solid-state electrolyte to construct sulfide solid-state batteries. LLO-Ru@S samples were obtained through Ru doping and surface modification. Detailed EIS, DRT, and XPS analyses revealed the underlying mechanisms of interface degradation. Through targeted improvements, a chemically / electrochemically compatible and stable interface was constructed, stabilizing lattice oxygen in the bulk phase and suppressing interface failure during the sulfidation process. For the LLO-Ru@S samples, Ru doping suppressed interface degradation caused by oxygen release and improved Li diffusion. The sulfide surface layer not only suppressed side reactions between the sulfide solid-state electrolyte and the cathode material under high voltage but also effectively mitigated the SCL effect, constructing interfacial charge transfer channels and significantly enhancing the initial activation of the Li2MnO3 phase. The strategy employed in this work opens a new avenue combining the advantages of sulfide solid-state electrolytes and lithium-rich manganese-based cathode materials, making it possible to develop suitable high-energy-density sulfide all-solid-state batteries.

[0096] It should be noted that the above description is only a preferred embodiment of the present invention and is not intended to limit the present invention. Any modifications, equivalent substitutions, improvements, etc., made within the spirit and principles of the present invention should be included within the protection scope of the present invention.

[0097] References:

[0098] [1] BWLiu,NF Hu,C.Li,J.Ma,JWZhang,Y.Yang,DYSun,BXYin,GLCui,Direct Observation of Li-Ion Transport Heterogeneity Induced by NanoscalePhase Separation in Li-rich Cathodes ofSolid-State Batteries.AngewandteChemie.2022,61,e202209626.

Claims

1. A lithium-rich manganese-based material, characterized in that, The lithium-rich manganese-based material has a core and a coating layer, wherein the core is Li 1+x TM 1-x Ru y O2, TM is selected from one or more of Mn, Ni and Co, 0<x<0.3, 0<y<0.02, the coating layer is a sulfur-containing coating layer, and the raw material of the sulfur-containing coating layer comprises CS2.

2. The lithium-rich manganese-based material according to claim 1, characterized in that, The Li 1+x TM 1-x Ru y The Ru element in O2 comes from one or more of Ru2O, ruthenium chloride, and ruthenium acetylacetonate.

3. The lithium-rich manganese-based material according to claim 1, characterized in that, The thickness of the sulfur-containing coating layer is 8nm-12nm.

4. The lithium-rich manganese-based material according to claim 1, characterized in that, The lithium-rich manganese-based material has a layered phase and a rock salt phase.

5. The lithium-rich manganese-based material according to claim 1, characterized in that, The thickness of the rock salt phase in the lithium-rich manganese-based material is not less than 2.5 nm.

6. The method for preparing the lithium-rich manganese-based material according to any one of claims 1-5, characterized in that, The preparation method includes: weighing lithium source, TM source and Ru source according to the elemental stoichiometry, grinding and mixing them evenly, and performing three-stage sintering to obtain lithium-rich manganese-based material precursor; adding CS2 and performing heat treatment to obtain the lithium-rich manganese-based material of the present invention.

7. The preparation method according to claim 6, characterized in that, The lithium source includes one or more of Li2CO3, lithium manganese oxide, lithium iron phosphate, lithium acetate, and lithium nitrate; the Ru source includes one or more of Ru2O, ruthenium chloride, and ruthenium acetylacetone.

8. The preparation method according to claim 6, characterized in that, The TM source includes one or more of Co, Mn, and Ni sources.

9. The preparation method according to claim 8, characterized in that, The Co source includes one or more of cobalt sulfate, cobalt nitrate, cobalt chloride, nickel acetate, and manganese-rich ternary carbonates.

10. The preparation method according to claim 8, characterized in that, The Mn source includes one or more of manganese sulfate, manganese nitrate, manganese chloride, manganese acetate, and manganese-rich ternary carbonates.

11. The preparation method according to claim 8, characterized in that, The Ni source includes one or more of nickel sulfate, nickel nitrate, nickel chloride, nickel acetate, and manganese-rich ternary carbonates.

12. The preparation method according to claim 6, characterized in that, The TM source is selected from the same type, including manganese-rich ternary carbonates.

13. The preparation method according to claim 12, characterized in that, The manganese-rich ternary carbonate includes (Ni) x Co y Mn 1-x-y One or more of CO3, wherein x>0, y>0, and x+y<1.

14. The preparation method according to claim 13, characterized in that, The (Ni) x Co y Mn 1-x-y CO3 includes Ni 0.16 Co 0.16 Mn 0.68 CO3.

15. The preparation method according to claim 6, characterized in that, The amount of lithium source added is 5% excess by mass.

16. The preparation method according to claim 6, characterized in that, The three-stage sintering process specifically includes a first stage with a sintering temperature of 300℃-400℃ and a sintering time of 4h-6h, a second stage with a sintering temperature of 680℃-720℃ and a sintering time of 14h-16h, and a third stage with a sintering temperature of 840℃-880℃ and a sintering time of 4h-6h.

17. The preparation method according to claim 6, characterized in that, The heating rate for the three-stage sintering is 2.5℃ / min-3.5℃ / min.

18. The preparation method according to claim 6, characterized in that, The heat treatment temperature is 150℃-250℃, and the heat treatment time is 0.8h-1.2h.

19. The preparation method according to claim 6, characterized in that, The conditions for the heat treatment also include being performed in a nitrogen atmosphere.

20. A cathode having the lithium-rich manganese-based material according to any one of claims 1-5.

21. An electrochemical device having the lithium-rich manganese-based material according to any one of claims 1-5.

22. The electrochemical device according to claim 21, characterized in that, The electrochemical device includes a battery.

23. The electrochemical device according to claim 22, characterized in that, The battery includes an all-solid-state ion battery.

24. The electrochemical device according to claim 23, characterized in that, The solid-state ion battery includes a sulfide all-solid-state electrolyte battery.

25. The electrochemical device according to claim 24, characterized in that, The sulfide all-solid-state electrolyte battery further includes a negative electrode and a sulfide solid-state electrolyte; the negative electrode is Li-In; and the sulfide all-solid-state electrolyte is Li6PS5Cl.

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