Preparation method and application of anthracite-based hard carbon sodium ion battery negative electrode material

Abstract

The application discloses a preparation method of anthracite-based hard carbon sodium ion battery negative electrode material and application thereof, and belongs to the field of battery negative electrode materials; the anthracite is used as raw material, and first, nitric acid washing and pre-oxidation are carried out to control the ash content and chemical composition of coal direct liquefaction residue; then, part of the silicon dioxide is removed by soaking in a hydrofluoric acid solution; and then, the anthracite-based hard carbon sodium ion battery negative electrode material is prepared through high-temperature carbonization; the anthracite-based hard carbon sodium ion battery negative electrode material prepared by the application has the advantages of high reversible specific capacity, low electrode reaction resistance, good rate performance, excellent cycle stability and high initial coulomb efficiency.

Description

Technical Field

[0001] This invention relates to the field of battery anode materials, and in particular to a method for preparing anthracite-based hard carbon sodium-ion battery anode material and its application. Background Technology

[0002] Lithium-ion batteries have been widely used in new energy vehicles and portable electronic devices. However, the further promotion of lithium-ion batteries in large-scale energy storage applications faces many challenges, including limited lithium resources, uneven geographical distribution, rising costs, and concerns about supply chain security. In contrast, sodium is extremely abundant and inexpensive in the Earth's crust and seawater, making sodium-ion batteries a promising alternative for large-scale energy storage and smart grid applications.

[0003] In sodium-ion batteries, the anode material remains one of the key factors limiting energy density and cycle life. In existing anode systems, commercial graphite is difficult to use due to its small interlayer spacing and the thermodynamic instability of sodium-graphite intercalation compounds, hindering the achievement of Na+ energy density and cycle life. + Reversible storage is possible. Although alloy and conversion anodes have high theoretical capacities, they typically suffer from significant volume changes, high operating potentials, and substantial capacity degradation after cycling. Hard carbon is currently considered one of the most promising anode materials for sodium-ion batteries due to its large interlayer spacing, low operating potential, excellent structural stability, abundant precursor supply, and cost-effectiveness. Sodium-ion storage in hard carbon is generally believed to follow a "ramp-plateau" coupling mechanism. The ramp region exhibits rapid reaction kinetics, and increasing the ramp capacity is beneficial for improving discharge rate and cycle stability. The plateau region has a lower discharge voltage, directly affecting the energy / power density and initial coulombic efficiency of sodium-ion batteries. However, increasing the capacity of the ramp region usually leads to a decrease in the initial coulombic efficiency in hard carbon. An excessively low initial coulombic efficiency promotes the formation of a thicker solid electrolyte interface layer, reducing conductivity and overall capacity. Furthermore, increasing closed porosity can improve plateau capacity but reduces tap density and compaction density, thus lowering volumetric energy density. The low pore-filling potential of closed pores also benefits sodium intercalation under high-rate conditions. Therefore, balancing ramp capacity and plateau capacity, and exploring the threshold of plateau capacity, is crucial for the performance of hard carbon anodes.

[0004] Coal-derived resources, especially anthracite, are characterized by high carbon content, high carbon yield, low cost, and high compaction density. Their rich aromatic structure allows for localized graphitization at high temperatures, making them ideal for producing low-cost, high-density hard carbon. However, the inherent mineral impurities (ash) in anthracite exert complex influences on microstructure and interfacial behavior during carbonization and electrochemical cycling. Traditional research typically aims to remove ash as thoroughly as possible (e.g., through acid washing) to suppress side reactions and improve initial coulombic efficiency. Recent studies have shown that under specific conditions, appropriate amounts of ash can promote localized graphitization, foster the formation of multi-scale porous structures, and actively participate in the formation and compositional regulation of the solid electrolyte interfacial film, thus positively impacting battery plateau capacity and cycle stability. These findings suggest that "controllable ash" has the potential to transform from a mere "harmful impurity" into a usable tool for microstructure and interfacial engineering. However, a systematic understanding of the synergistic relationship between "ash-carbon framework-pore structure-interfacial chemistry" remains insufficient in current coal-based hard carbon research. Summary of the Invention

[0005] The purpose of this invention is to provide a method for preparing anthracite-based hard carbon sodium-ion battery anode material and its application, using anthracite as raw material to prepare a hard carbon sodium-ion battery anode material.

[0006] The technical solution adopted in this invention is as follows:

[0007] In a first aspect, the present invention provides a method for preparing anthracite-based hard carbon sodium-ion battery anode material, comprising:

[0008] Step 1: Soak anthracite powder in HNO3 solution at room temperature to remove some metal elements and introduce some oxygen-containing functional groups, including hydroxyl, carboxyl, and carbonyl groups. After soaking in HNO3 solution, filter out the solid residue, wash the residue with distilled water until neutral, and dry to obtain metal-free anthracite powder. The concentration of HNO3 solution used is 1~3 mol / L, and the anthracite powder and HNO3 solution are prepared at a mass-volume ratio of 1g:4~8mL. The soaking time is 4~8h.

[0009] Step 2: Soak the anthracite powder obtained in Step 1 in hydrofluoric acid solution at a temperature of 30-70℃ to remove some of the silica. After the hydrofluoric acid solution soaking is completed, filter out the residue, wash the residue with distilled water until neutral, and dry to obtain the silica-free anthracite powder. The concentration of the hydrofluoric acid solution used is 1-3 mol / L, and the anthracite powder and hydrofluoric acid solution are prepared at a mass-volume ratio of 1g:4-8mL. The soaking time is 4-8h.

[0010] Step 3: Heat and carbonize the desiliconized anthracite powder in a tube furnace; under a nitrogen atmosphere, heat to 1100~1500 ℃ at a heating rate of 5℃ / min and hold at this temperature for 0.5~4 h; then cool naturally to room temperature to obtain anthracite-based hard carbon sodium-ion battery anode material.

[0011] Preferably, anthracite powder is taken from powder with a mesh size between 200 and 250 as raw material.

[0012] In a preferred embodiment, in step 1, the concentration of the HNO3 solution used is 2 mol / L, and the anthracite powder and HNO3 solution are prepared at a mass-volume ratio of 1 g: 6 mL, with a soaking time of 6 h; in step 2, the anthracite powder obtained in step 1 is soaked in hydrofluoric acid solution at a temperature of 50°C; the concentration of the hydrofluoric acid solution used is 2 mol / L, and the anthracite powder and hydrofluoric acid solution are prepared at a mass-volume ratio of 1 g: 6 mL, with a soaking time of 6 h; in step 3, the anthracite powder is heated to 1300°C at a heating rate of 5°C / min under a nitrogen atmosphere and held at this temperature for 2 h; then it is naturally cooled to room temperature to obtain the anthracite-based hard carbon sodium-ion battery anode material.

[0013] Secondly, the present invention provides an application of anthracite-based hard carbon sodium-ion battery anode material, wherein the anthracite-based hard carbon sodium-ion battery anode material prepared by the aforementioned preparation method is used as a sodium-ion battery anode material.

[0014] Preferably, the working electrode is prepared by the following method: the working electrode is prepared by mixing 80 wt.% hard carbon, 10 wt.% Super-P carbon black and 10 wt.% sodium alginate in an aqueous solution; the mass loading of the electrode is 1.8~2.6 mg / cm³. 2 Whatman glass fiber is used as the diaphragm, and 1M NaPF6 dissolved in ethylene glycol dimethyl ether (100 vol%) is used as the electrolyte. All components are vacuum dried for 12 hours to remove moisture. The gas in the glove box must be replaced more than three times before operation to ensure an inert atmosphere. The water and oxygen content in the glove box is <1 ppm, the vacuum drying oven is 60-80℃, and the tablet press pressure is 50 MPa. The assembly sequence is: positive electrode shell first, positive electrode shell → positive electrode plate → electrolyte → diaphragm → sodium plate → gasket → spring plate → negative electrode shell. The electrolyte injection error must be controlled within ±0.02 mL, and the diaphragm must be completely impregnated. A digital display tablet press is used to seal at a pressure of 50 MPa to ensure a contact resistance <0.5 mΩ.

[0015] The beneficial effects of this invention are as follows: This invention provides a method for preparing anthracite-based hard carbon sodium-ion battery anode material and its application. Using anthracite as raw material, the invention first pre-oxidizes the material by nitric acid elution to control the ash content and chemical composition of the coal direct liquefaction residue; then, it removes some silica by soaking in hydrofluoric acid solution; finally, it prepares the anthracite-based hard carbon sodium-ion battery anode material through high-temperature carbonization. The anthracite-based hard carbon sodium-ion battery anode material prepared by this invention possesses high reversible specific capacity, low electrode reaction resistance, good rate performance, excellent cycle stability, and high initial coulombic efficiency.

[0016] Finally, the obtained coal-based hard carbon exhibited a capacity of 348.5 mAh / g (capacity retention of 93.4%) after 250 cycles at a current density of 0.1 A / g, with an initial coulombic efficiency of 85.1%, and maintained 186.5 mAh / g (capacity retention of 90.6%) after 1000 cycles at a current density of 1.0 A / g. Furthermore, a full cell assembled using the obtained hard carbon anode (NVP / / CA1300) showed an initial discharge capacity of 186.6 mAh / g at 0.05 A / g, and maintained 84.6 mAh / g at 1.0 A / g after 900 cycles, with a coulombic efficiency approaching 100%. These results validate the practical application potential of anthracite-derived hard carbon in real-world solid-state electrolyte battery configurations and provide a structural paradigm for the synergistic design of coal-based hard carbon anodes through "porosity-long-range order-short-range disorder-ash" engineering. Attached Figure Description

[0017] Figure 1 The diagram shows the preparation process and structural evolution of the carbon anodes prepared from anthracite at different carbonization temperatures according to the present invention.

[0018] Figure 2 The images shown are scanning electron microscope images of CA1100, CA1300, and CA1500.

[0019] Figure 3 The image shows the elemental spectra of CA1100, CA130, and CA1500.

[0020] Figure 4 This is a graph showing the elemental content of calcium, magnesium, aluminum, and silicon in three samples (CA1100, CA130, and CA1500) obtained through energy dispersive spectroscopy (EDS) analysis.

[0021] Figure 5 The figures show the N2 adsorption-desorption isotherms and pore size distribution of CA1100, CA1300 and CA1500.

[0022] Figure 6These are the small-angle scattering spectra of CA1100 and CA1300.

[0023] Figure 7 These are the TGA curves for CA1100, CA1300, and CA1500.

[0024] Figure 8 These are the XRD patterns of CA1100, CA1300, and CA1500.

[0025] Figure 9 These are XPS spectra of CA1100, CA1300, and CA1500.

[0026] Figure 10 These are high-resolution transmission electron microscope images from the CA1100.

[0027] Figure 11 These are high-resolution transmission electron microscope images from a CA1300.

[0028] Figure 12 These are high-resolution transmission electron microscope images from a CA1500.

[0029] Figure 13 These are the XRD fitted spectra of CA1100, CA1300, and CA1500.

[0030] Figure 14 These are the fitted Raman spectra of CA1100, CA1300, and CA1500.

[0031] Figure 15 These are high-resolution C1sXPS spectra of CA1100, CA1300, and CA1500.

[0032] Figure 16 These are the performance of CA1100, CA1300 and CA1500 electrodes at different current densities and the charge-discharge curves of CA1300 at different voltages.

[0033] Figure 17 The cycling performance of CA1100, CA1300 and CA1500 electrodes was tested at a current of 0.1 A / g; and the charge-discharge curves of the CA1300 electrode were obtained at different cycles under this current.

[0034] Figure 18 These are the initial coulombic efficiency values ​​of the CA1100, CA1300, and CA1500 electrodes at a current of 0.1 A / g.

[0035] Figure 19 The first cycle charge-discharge curves of different samples at 0.1 A / g are shown, along with a comparison of the corresponding plateau capacity and ramp capacity.

[0036] Figure 20 The capacity retention rate of CA1100, CA1300 and CA1500 electrodes after 250 cycles at a current of 0.1 A / g.

[0037] Figure 21 The cycling performance of CA1100, CA1300 and CA1500 electrodes at 1.0 A / g is shown.

[0038] Figure 22 These are the in-situ EIS spectra of the CA1100, CA1300, and CA1500 electrodes during their first cycle, along with the corresponding analog circuit diagrams.

[0039] Figure 23 It is the AC impedance fitting resistor for the CA1300 electrode.

[0040] Figure 24 It is the AC impedance fitting resistor for the CA1100 electrode.

[0041] Figure 25 It is the AC impedance fitting resistor for the CA1100 electrode.

[0042] Figure 26 These are in-situ EIS spectra of CA1100, CA1300, and CA1500 electrodes during their first cycle.

[0043] Figure 27 These are the GITT curves for the CA1100, CA1300, and CA1500 electrodes.

[0044] Figure 28 Na obtained through GITT + Diffusion coefficient (logD).

[0045] Figure 29 These are the CV curves of CA1100, CA1300, and CA1500 samples at scan rates of 0.1–1.0 mV / s.

[0046] Figure 30 This is a graph showing the contribution ratio of capacitance to diffusion control for CA1100, CA1300, and CA1500 at different scan rates.

[0047] Figure 31 This is a capacitive contribution analysis graph of CA1100, CA1300, and CA1500 at a scan rate of 0.6 mV / s.

[0048] Figure 32 This is the CV curve of NVP / / CA1300 total sodium ions.

[0049] Figure 33This is a graph showing the rate performance test results of the NVP / / CA1300 all-sodium ion battery at different current densities.

[0050] Figure 34 These are the charge-discharge curves of the NVP / / CA1300 all-sodium ion battery at different current densities.

[0051] Figure 35 These are the cycle stability and coulombic efficiency at 1.0 A / g for 900 cycles. Detailed Implementation

[0052] The technical solutions in the embodiments of the present invention will be clearly and completely described below. Obviously, the described embodiments are only some embodiments of the present invention, and not all embodiments. Based on the embodiments of the present invention, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of the present invention.

[0053] This invention uses anthracite coal provided by Inner Mongolia Shanghai Temple Mining Co., Ltd. as raw material. The main components and content of the anthracite coal are shown in Table 1 below. The anthracite coal is pulverized, and powder passing through a mesh between 200 and 250 is used as the raw material in this invention.

[0054] Table 1. Statistical table of main components and content of anthracite.

[0055] Example 1

[0056] Step 1: Soak anthracite powder in HNO3 solution at room temperature to remove some metal elements and introduce some oxygen-containing functional groups (hydroxyl, carboxyl, carbonyl). After the HNO3 solution soaking treatment is completed, filter out the solid filter residue, then wash the filter residue with distilled water until neutral and dry to obtain metal-free anthracite powder.

[0057] In this embodiment, a 2 mol / L HNO3 solution was used. The anthracite powder and HNO3 solution were prepared at a mass-to-volume ratio of 1 g: 6 mL, and the soaking time was 6 hours. More specifically, in this embodiment, the amount of anthracite powder used was 10 g, and the amount of HNO3 solution used was 60 mL.

[0058] Step 2: The anthracite powder obtained in Step 1 is soaked in hydrofluoric acid solution at a temperature of 30-70℃ to remove some of the silica. After the hydrofluoric acid soaking treatment is completed, the filter residue is filtered out, then washed with distilled water until neutral and dried to obtain silica-free anthracite powder.

[0059] In this embodiment, a 2 mol / L hydrofluoric acid solution was used. The anthracite powder (excluding metal) was mixed with the hydrofluoric acid solution at a mass-to-volume ratio of 1 g:6 mL. The temperature was 50°C, and the soaking time was 6 hours. More specifically, in this embodiment, the amount of anthracite powder (excluding metal) was 10 g, and the amount of hydrofluoric acid solution was 60 mL.

[0060] Step 3: Place 2g of desiliconized anthracite powder into a container, such as a ceramic boat; then place the container into a tube furnace. Under a nitrogen atmosphere, heat to 1300℃ at a heating rate of 5℃ / min and maintain this temperature for 2 hours. Then allow it to cool naturally to room temperature to obtain the anthracite-based hard carbon sodium-ion battery anode material.

[0061] The sample obtained in this embodiment is named CA1300 according to its carbonization temperature. Example 2

[0062] Step 1: At room temperature, soak anthracite powder in a 3 mol / L HNO3 solution. In this embodiment, the anthracite powder and HNO3 solution are prepared at a mass-to-volume ratio of 1 g:8 mL, and the soaking time is 8 hours. More specifically, in this embodiment, the amount of anthracite powder used is 10 g, and the amount of HNO3 solution used is 80 mL. After the HNO3 solution soaking treatment is completed, the solid filter residue is filtered out, then washed with distilled water until neutral and dried to obtain metal-free anthracite powder.

[0063] Step 2: The anthracite powder obtained in Step 1, after soaking in a 3 mol / L hydrofluoric acid solution at 70°C, is partially soaked to remove silica. In this embodiment, the anthracite powder and hydrofluoric acid solution are prepared at a mass-to-volume ratio of 1 g:8 mL, and the soaking time is 8 hours. More specifically, in this embodiment, the amount of anthracite powder used is 10 g, and the amount of hydrofluoric acid solution used is 80 mL. After the hydrofluoric acid solution soaking treatment is completed, the filter residue is filtered out, then washed with distilled water until neutral and dried to obtain silica-free anthracite powder.

[0064] Step 3: Place 2g of desiliconized anthracite powder into a ceramic boat, and then place the container into a tube furnace. Under a nitrogen atmosphere, heat to 1500℃ at a heating rate of 5℃ / min, and maintain this temperature for 4 hours. Then allow it to cool naturally to room temperature to obtain anthracite-based hard carbon sodium-ion battery anode material.

[0065] The sample obtained in this embodiment is named CA1500 according to its carbonization temperature. Example 3

[0066] Step 1: At room temperature, soak the anthracite powder in a 1 mol / L HNO3 solution. In this embodiment, the anthracite powder and HNO3 solution are prepared at a mass-to-volume ratio of 1 g:4 mL, and the soaking time is 4 hours. More specifically, in this embodiment, the amount of anthracite powder used is 10 g, and the amount of HNO3 solution used is 40 mL. After the HNO3 solution soaking treatment is completed, the solid filter residue is filtered out, then washed with distilled water until neutral and dried to obtain metal-free anthracite powder.

[0067] Step 2: The anthracite powder obtained in Step 1, containing removed metals, is soaked in a 1 mol / L hydrofluoric acid solution at 30°C to remove some of the silica. In this embodiment, the anthracite powder containing removed metals is prepared with hydrofluoric acid solution at a mass-to-volume ratio of 1 g:4 mL, and the soaking time is 4 hours. More specifically, in this embodiment, the amount of anthracite powder containing removed metals is 10 g, and the amount of hydrofluoric acid solution is 40 mL. After the hydrofluoric acid solution soaking treatment is completed, the filter residue is filtered out, then washed with distilled water until neutral and dried to obtain silica-free anthracite powder.

[0068] Step 3: Place 2g of desiliconized anthracite powder into a ceramic boat, and then place the container into a tube furnace. Under a nitrogen atmosphere, heat to 1100℃ at a heating rate of 5℃ / min, and maintain this temperature for 0.5h. Then allow it to cool naturally to room temperature to obtain anthracite-based hard carbon sodium-ion battery anode material.

[0069] The sample obtained in this embodiment is named CA1100 according to its carbonization temperature.

[0070] Comparative Example 1

[0071] The HNO3 solution used in step 1 of Example 1 was replaced with an HCl solution to treat the anthracite powder. The amount and concentration of the HCl solution used were the same as those used in the HNO3 solution in Example 1. The anthracite-based hard carbon obtained in Comparative Example 1 was named CA1300HCl. The specific steps are as follows.

[0072] Step 1: Soak the anthracite powder in HCl solution at room temperature. After the HCl solution soaking is completed, filter out the solid residue, then wash the residue with distilled water until neutral and dry to obtain metal-free anthracite powder.

[0073] In this embodiment, a 2 mol / L HCl solution was used, and the anthracite powder was mixed with the HCl solution at a mass-to-volume ratio of 1 g: 6 mL, with a soaking time of 6 hours. More specifically, in this embodiment, the amount of anthracite powder used was 10 g, and the amount of HCl solution used was 60 mL.

[0074] Steps 2 and 3 are the same as in Example 1.

[0075] Comparative Example 2

[0076] Perform steps 1 and 3 of Example 1, but not step 2. The anthracite-based hard carbon obtained in Comparative Example 2 is named CA1300SiO.

[0077] This invention employs solid-state carbon NMR spectroscopy to analyze the acid-treated anthracite intermediates from Examples 1-3 and Comparative Examples 1 and 2, investigating their oxygen-containing group content. The results are shown in Table 2 below. In Table 2, the units for hydroxyl, carboxyl, and carbonyl groups are mmol / g.

[0078] Table 2. Analysis of oxygen-containing group content in the intermediate products of anthracite after acid treatment during the preparation of Examples 1-3 and Comparative Examples 1 and 2.

[0079]

[0080] The comparative analysis is as follows: In Example 1, step 1 used a nitric acid solution, while in Comparative Example 1, step 1 used an HCl solution. Comparing the content of oxygen-containing groups (hydroxyl, carboxyl, carbonyl) in the products after treatment, the data in Table 2 shows that the amount of oxygen-containing groups in the product treated with nitric acid solution is significantly higher than that in the product treated with hydrochloric acid. A comparison of Examples 1, 2, and 3 reveals that the amount of oxygen-containing groups increases with increasing nitric acid concentration, dosage, and treatment time.

[0081] The present invention performs elemental analysis on the anthracite intermediate products after acid treatment during the preparation process of the examples and comparative examples. The results are shown in Table 3 below. Comparing Example 1 and Comparative Example 1, it was found that the ash content in the anthracite intermediate products was similar when using nitric acid and hydrochloric acid as detergents. Comparing Example 1 and Comparative Example 2, it was found that the silicon content in the anthracite intermediate product without hydrofluoric acid treatment was significantly higher than that in the anthracite intermediate product treated with hydrofluoric acid.

[0082] Table 3 shows the analysis of ash and silicon content in the acid-treated anthracite intermediates during the preparation of Examples 1-3 and Comparative Examples 1 and 2.

[0083]

[0084] This invention employs Raman spectroscopy to analyze the hard carbon anode materials prepared in the examples and comparative examples. The results show that all hard carbon materials prepared in this invention exhibit Raman spectroscopy values ​​at approximately 1350 and 1580 cm⁻¹. -1 All exhibit typical D and G bands, corresponding to the breathing vibration of disordered or defect-activated sp2 hybrid carbon rings and the in-plane stretching vibration of sp2 hybrid carbon atoms, respectively.

[0085] The intensity ratio of D-band to G-band (I) D / I GIt is typically used to assess the degree of defect and graphitization level of carbon materials; D / I G The higher the value, the lower the level of graphitization. This invention calculated the strength ratio (Istrength) of the hard carbon anode materials prepared in the examples and comparative examples. D / I G As shown in Table 4.

[0086] Table 4. Strength ratio of hard carbon anode materials prepared in the examples and comparative examples (I) D / I G Analysis table.

[0087]

[0088] The results show that the introduction of oxygen-containing functional groups such as hydroxyl and carboxyl groups creates steric hindrance between aromatic rings, inhibiting the directional stacking of aromatic units during carbonization and reducing the degree of graphitization of the material. Furthermore, alkali metals (such as K and Na), alkaline earth metals (such as Ca and Mg), and transition metals (such as Fe) present in the ash of anthracite intermediates exhibit significant catalytic graphitization effects during carbonization. When the carbonization temperature reaches above 550℃, these minerals act as catalysts, promoting the rearrangement and microcrystal growth of the carbon hexagonal network structure, accelerating the transformation of coal-based carbon materials into graphite-like structures. For hard carbon materials, the higher the carbonization temperature and the longer the carbonization time, the higher the degree of graphitization. Comparatively, the acid-treated anthracite intermediate prepared in Example 2 contained the most oxygen-containing functional groups and the least ash, and should have the lowest level of graphitization. However, Example 2 had the highest carbonization temperature and the longest carbonization time, resulting in an increased level of hard carbon graphitization. Therefore, under the combined effects of factors such as oxygen-containing groups, ash content, carbonization temperature, and carbonization time, the hard carbon obtained in Example 1 exhibited the lowest degree of graphitization. Except for the difference in oxygen-containing group content, Comparative Example 1 and Example 1 had similar ash contents, identical carbonization temperatures, and carbonization times. However, Comparative Example 1 showed a significantly higher level of graphitization than Example 1, further demonstrating that a high oxygen-containing group content is beneficial for suppressing excessive graphitization of hard carbon materials. Conversely, except for the difference in ash content, Comparative Example 2 and Example 1 had similar oxygen-containing group contents, identical carbonization temperatures, and carbonization times. However, Comparative Example 2 showed a significantly higher level of graphitization than Example 1, again indicating that excessively high ash content is beneficial for accelerating the graphitization of hard carbon materials.

[0089] In this invention, working electrodes were prepared from samples of Examples 1-3 and Comparative Examples 1-2, and their electrochemical performance was tested. The specific methods are as follows:

[0090] The working electrode is prepared by mixing 80 wt.% hard carbon, 10 wt.% Super-P carbon black, and 10 wt.% sodium alginate in an aqueous solution. The electrode's mass loading is 1.8~2.6 mg / cm³. 2Whatman glass fiber (GF / D) was used as the separator, and 1M NaPF6 dissolved in dimethyl ethylene glycol (DME) = 100 vol% was used as the electrolyte.

[0091] Sodium electrode (counter electrode), separator, electrolyte (dimethyl ether solution containing NaPF6). All components (battery casing, electrodes, separator, etc.) must be vacuum dried for 12 hours to remove moisture. The gas must be purged at least three times before operation in the glove box to ensure an inert atmosphere. Equipment includes: glove box (water and oxygen content <1ppm), vacuum drying oven (60-80℃), tablet press (pressure 50MPa), and pipette (for precise control of electrolyte volume).

[0092] Assembly sequence: (starting with positive electrode shell) Positive electrode shell → positive electrode plate → electrolyte (0.1-0.2mL) → separator (wetted with electrolyte) → sodium plate (negative electrode) → gasket → spring contact → negative electrode shell. The electrolyte injection error must be controlled within ±0.02mL, and the separator must be completely soaked. Use a digital display tablet press to seal at 50MPa pressure, ensuring contact resistance <0.5mΩ. After standing for 4 hours, check the open circuit voltage. Abnormal values ​​(e.g., <2.5V) require investigation of short circuit risk.

[0093] Electrochemical experiments were conducted using 2025 coin cells, and the testing methods are as follows:

[0094] At room temperature, within a voltage range of 0.01~3.0V (relative to Na / Na) + The hard carbon electrode was subjected to cycle performance testing at a current density of 100 mA / g on the Land testing system. The test results are shown in Table 5 below.

[0095] Table 5. Summary of electrochemical performance test results of hard carbon anodes prepared in the examples and comparative examples.

[0096]

[0097] Comparative analysis revealed that with increasing graphitization, the initial discharge specific capacity and initial coulombic efficiency of hard carbon decreased. This is because the molecular structure of coal is rich in polycyclic aromatic hydrocarbon units. During high-temperature carbonization, these aromatic units undergo directional stacking due to strong intermolecular π-π conjugation. This stacking effect gradually eliminates stacking faults in the carbon layers, and adjacent hexagonal planar network layers of carbon exhibit an ordered arrangement, ultimately forming a soft carbon-like structure with smaller interlayer spacing (increased graphitization).

[0098] From the perspective of sodium storage mechanisms, sodium ion storage in carbon materials mainly relies on two methods: defect adsorption and interlayer intercalation. When the interlayer spacing of coal-based carbon materials narrows to a certain extent, the interlayer intercalation channels for sodium ions are significantly compressed. On the one hand, sodium ions with relatively large radii (approximately 0.102 nm) struggle to overcome the steric hindrance effect between layers to complete effective insertion; on the other hand, even if some sodium ions successfully insert, they are bound by interlayer van der Waals forces during extraction, hindering the insertion-extraction kinetics. Furthermore, narrowing the interlayer spacing reduces defect sites on the material surface, weakening the adsorption and storage capacity of sodium ions. This causes the sodium storage mechanism to gradually shift towards a single interlayer intercalation mechanism, ultimately resulting in a smaller reversible sodium storage capacity, lower initial coulombic efficiency, and reduced capacity retention, making it difficult to meet the needs of commercial applications. Additionally, silicon in hard carbon materials mainly exists in the form of silicon oxide, which reacts to form Na during the initial sodium insertion process. x SiO x And Na2O, which reduces the initial coulombic efficiency.

[0099] To further analyze the effects of carbonization temperature, carbonization time, and ash content on the graphitization degree of hard carbon materials, the present invention conducted detailed physicochemical characterization and electrochemical performance testing on the hard carbon materials of Examples 1-3.

[0100] like Figure 1 The diagram illustrates the preparation process and structural evolution of carbon anodes prepared from anthracite at different carbonization temperatures according to this invention. By precisely controlling the calcination temperature and holding time, the obtained hard carbon exhibits a microstructure characterized by the coexistence of long-range order and short-range disorder. Based on the different calcination temperatures, these samples are named CA1100, CA1300, and CA1500, respectively, corresponding to 1100... o C, 1300 o C and 1500 o The calcination temperature of C.

[0101] To elucidate the structural evolution of anthracite during carbonization, scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDS) were used to systematically analyze samples prepared at different calcination temperatures. Figure 2 The images shown are scanning electron microscope (SEM) images of CA1100, CA1300, and CA1500. Image a on the left is an SEM image of CA1100; image b in the middle is an SEM image of CA1300; and image c on the right is an SEM image of CA1500.

[0102] Scanning electron microscope images of CA1100 ( Figure 2 As shown in a), its surface is covered with a large number of fine and uniformly distributed spherical ash particles, which is closely related to mineral precipitation and organic matter decomposition during pyrolysis. When the temperature rises to 1300...o C, when the carbonization time reaches 2 hours ( Figure 2 (b) The number of ash particles on the surface decreased significantly, while their size increased, indicating that the ash underwent migration, aggregation, and coarsening under the influence of temperature. Further heating to 1500... o C, after the time is extended to 4 hours ( Figure 2 c) Most of the ash melts and permeates into the pores or escapes through dripping / exudation, resulting in a surface that is virtually "ash-free." A small amount of uniformly dispersed ash modulates the chemistry of the solid electrolyte interphase (SEI) membrane during the initial electrolyte decomposition process, forming a denser, inorganic-rich SEI membrane, thereby reducing interfacial impedance and suppressing side reactions. Temperature-driven ash evolution alters particle morphology and modulates interfacial chemistry, thereby improving cycle stability and rate performance.

[0103] Figure 3 The image shows the energy spectrum elements of CA1100, CA130, and CA1500. The first row (a) is the energy spectrum element diagram of CA1100, the second row (b) is the energy spectrum element diagram of CA1300, and the third row (c) is the energy spectrum element diagram of CA1500.

[0104] Figure 3 The EDS elemental distribution results shown in Figure ac indicate that the carbon framework and inorganic ash components exhibit a synergistic evolution in anthracite-derived hard carbon treated at different carbonization temperatures. In all three samples, carbon is uniformly distributed within the particles, forming a continuous matrix regardless of carbonization conditions, which facilitates efficient electron transport. In contrast, oxygen, calcium, magnesium, aluminum, and silicon are locally enriched, indicating that inorganic phases from the coal precursor still exist as residual ash. These phases are mainly composed of aluminosilicates (Al-Si-O) and alkaline earth minerals (Ca, Mg), typical of coal-derived carbon components.

[0105] like Figure 4 This is a graph showing the elemental content of calcium, magnesium, aluminum, and silicon in three samples (CA1100, CA130, and CA1500) obtained through energy dispersive spectroscopy (EDS) analysis. Figure a shows the calcium content of the three samples, Figure b shows the magnesium content, Figure c shows the aluminum content, and Figure d shows the aluminum content.

[0106] Table 6. Elemental analysis results (wt.%) of CA1100, CA1300, and CA1500.

[0107]

[0108] according to Figure 4Table 6 shows significant differences between samples taken at different temperatures. In the CA1100 sample, the strong and heterogeneous signals of Ca, Mg, Al, and Si indicate incomplete thermal remodeling at low temperatures, with the presence of coarse-grained or aggregated mineral residues. The high oxygen content suggests that oxygen groups and oxides were not completely removed. In the CA1300 sample, the weaker and more uniform inorganic signals indicate a more uniform ash distribution. These dispersed materials act as in-situ rigid templates during carbon layer rearrangement and pore evolution, preventing structural collapse and maintaining a hierarchical pore structure. They also promote local carbon ordering, contributing to the formation of a continuous sp2 carbon network while preserving short-range disorder and defects, consistent with previous findings. Increasing the carbonization temperature to 1500... o Carbon (C) significantly reduces the signal intensity of non-carbon inorganic elements, indicating that ash components melt, volatilize, or are encapsulated within the carbon matrix, thus reducing detectable ash content. While higher temperatures enhance the ordered arrangement of carbon, excessive ash loss disrupts the structural support of the porous framework, leading to a denser carbon structure. Oxygen is present in all samples, and its content decreases with increasing temperature. Residual oxygen can reside in stable functional groups or inorganic oxides, which can improve electrolyte wettability and the reversibility of interfacial reactions.

[0109] To further investigate the evolution of the pore structure of anthracite at different temperatures, N2 adsorption-desorption tests were performed on the samples. Figure 5 The figures show the N2 adsorption-desorption isotherms and pore size distribution of CA1100, CA1300 and CA1500.

[0110] The results show that all three groups of samples exhibit typical type IV isotherms with H4 hysteresis loops, indicating that their pore networks are mainly composed of slit-like micropores / mesopores. CA1100 shows a relatively narrow pore volume distribution in the 2–20 nm range, indicating a simple and uniform pore structure. In contrast, CA1300 exhibits two distinct peaks in the 3–5 nm and 10–30 nm ranges, suggesting a hierarchical pore structure integrating mesopores and macropores. However, for CA1500, the proportion of small to medium-sized mesopores decreases significantly, while the proportion of pores larger than 20 nm increases, implying that under high-temperature conditions, some pore walls underwent structural rearrangement or local collapse.

[0111] Table 7 shows the pore structure parameters of CA1100, CA1300, and CA1500 composite materials.

[0112]

[0113] The quantitative data on surface area and pore volume presented in Table 7 show that although the surface area of ​​CA1300 is slightly lower than that of CA1100, its total pore volume and average pore size are significantly increased. This improves the balance between moderate surface area and efficient layered pore structure, which is beneficial for electrolyte wetting and sodium ion diffusion, while reducing the risk of serious side reactions caused by excessive surface area.

[0114] Small-angle X-ray scattering (SAXS) further confirmed the moderating effect of pore structure. For example... Figure 6 These are the small-angle scattering spectra of CA1100 and CA1300. From Figure 6 As can be seen, CA1100 exhibits higher scattering intensity at low q values, indicating a higher proportion of open pores and a relatively loose overall framework. In contrast, CA1300 shows a more significant decrease in scattering intensity in the medium q region, with a slight "inflection point," which is typically associated with an increase in closed pores. Based on the N2 adsorption-desorption results, it can be inferred that when the temperature increases from 1100... o C rose to 1300 o At C, some open pores transform into closed pores ranging from sub-nanometer to several nanometer scale. These closed pores are difficult to detect with N₂ at 77 K, but they are detected at the low potential plateau of Na in hard carbon. + It plays a crucial role in storage.

[0115] Therefore, while CA1300 retains a suitable number of interconnected mesopores to facilitate rapid mass transport, it significantly increases the number of closed pores, thereby contributing to plateau capacity and forming a pore structure that balances both mass transport and energy storage. In contrast, CA1100 has limited plateau capacity, while CA1500 suffers from decreased Na storage activity due to excessive compaction and collapse of some closed pores.

[0116] Thermogravimetric analysis (TGA) revealed the effect of temperature on ash content. Figure 7 This is a TGA curve graph for CA1100, CA1300, and CA1500. The graph shows that as the carbonization temperature increases from 1100... o C rises to 1500 o C, the weight loss curve shifts towards higher temperatures and the slope becomes steeper, indicating that volatiles and oxygen-containing groups are removed, while the degree of carbonization also increases. At 800 oAt temperature C, the residual mass of CA1100, CA1300, and CA1500 are approximately 5.4%, 2.4%, and 0.5%, respectively, mainly derived from refractory minerals. The high ash content provides structural support but increases inert matter and clogs pores. CA1500 products are almost free of impurities, but due to the lack of an inorganic framework structure, they are prone to interlayer shrinkage and porosity degradation at high temperatures.

[0117] The evolution of structural order was further confirmed by XRD results. Figure 8 These are the XRD patterns of CA1100, CA1300, and CA1500. As can be seen from the figures, all three samples exhibit the broad (002) diffraction peak characteristic of hard carbon, appearing at 26°; simultaneously, a diffraction signal related to planar order appears at approximately 43°, indicating that the material mainly exhibits a vortex carbon structure of "short-range order / long-range disorder." With the calcination temperature increasing from 1100... o C increased to 1300 o The C (002) peak intensity increased and became sharper, with the peak at 43° becoming clearer, indicating increased carbon layer stacking, improved microcrystalline crystallinity, and larger graphitic structure. However, when the temperature was further increased to 1500... o At temperature C, the intensity of peaks (002) and (101) weakens and broadens, showing a trend of first increasing and then decreasing. This indicates that high temperature induces the shrinkage and reorganization of the carbon skeleton, thereby causing bending or dislocation of the carbon layer and changes in the pore structure, which in turn reduces the consistency of XRD and the orderliness of the structure. CA1300 achieves a more reasonable balance between locally ordered structure and disordered / porous structure, and therefore can be considered closer to the ideal structural state required for high-performance hard carbon anodes.

[0118] XPS analysis revealed the evolution of surface functional groups. Figure 9 Figure 1 shows the XPS spectra of CA1100, CA1300, and CA1500, where Figure 1a is the XPS spectrum and Figure 1b is the O1s XPS spectrum. The complete spectra show that all samples are mainly composed of carbon and oxygen, with the oxygen content decreasing significantly with increasing calcination conditions. The high-resolution O1s spectrum can be decomposed into peaks at approximately 530.6 eV (C=O), 531.9 eV (CO), and 536.1 eV (OC=O). The intensity of these peaks decreases significantly with increasing calcination conditions, especially the C=O peak. This indicates that the active oxygen-containing functional groups are effectively decomposed at higher temperatures and longer calcination times, resulting in a cleaner carbon surface.

[0119] The O1s spectrum of CA1300 shows that although most of the reactive oxygen groups have been removed, some polar sites remain. These residual sites help improve the wettability of the electrolyte and promote the formation of a uniform solid electrolyte film. In contrast, the oxygen peak of CA1100 is more pronounced, indicating that in the first cycle, excessive polar functional groups promote the formation of a thicker solid electrolyte film, thereby consuming a large amount of sodium ions and increasing the occurrence of side reactions. Meanwhile, in the O1s spectrum of CA1500, the oxygen peak almost disappears completely, indicating that its surface is extremely oxygen-deficient. While this condition is beneficial for suppressing side reactions, it also reduces the polar sites and surface active centers that interact with the electrolyte.

[0120] Figure 10 These are high-resolution transmission electron microscope (TEM) images of the CA1100. Figure a is a high-resolution TEM image of the CA1100 sample, while figures b and c are high-magnification magnified views.

[0121] Figure 11 These are high-resolution transmission electron microscope (TEM) images of the CA1300. Figure a is a high-resolution TEM image of the CA1300 sample, while figures b and c are high-magnification magnified views.

[0122] Figure 12 These are high-resolution transmission electron microscope (TEM) images of the CA1500. Figure a is a high-resolution TEM image of the CA1500 sample, while figures b and c are high-magnification magnified views.

[0123] from Figures 10-12 As can be seen from this, as the temperature increases from 1100... o C rises to 1500 o C. The carbon layer structure transforms from a relatively regular stacking to an increasingly curved and fluctuating short-range disordered state. In CA1100, the carbon layers are tightly stacked, with a local interlayer spacing of approximately 0.344 nm, exhibiting typical graphite properties. This is beneficial for electron transport but hinders efficient sodium ion insertion. Figure 10 In CA1300, ordered microcrystals and disordered regions coexist, with localized areas exhibiting a multi-scale porous structure and an average interlayer spacing extending to approximately 0.357 nm. Furthermore, the (002) and (101) rings in the SAED pattern indicate the formation of graphite-like microcrystals with well-defined length scales at this temperature. This structure achieves a good balance between electronic conductivity and sodium ion diffusion, providing more readily available and efficient diffusion channels, thus ensuring an optimal balance between structural stability and sodium storage activity. Figure 11 When the temperature rises to 1500 oAt temperature C, the fringes of the carbon layer become more curved, looser, and more volatile, with a local interlayer spacing reaching 0.417 nm. This indicates significant structural reorganization and interlayer stacking relaxation at higher temperatures. While this provides more space for sodium ion insertion, it sacrifices some of the framework's rigidity and electronic continuity. Figure 12 ).

[0124] Figure 13 These are the XRD fitted spectra of CA1100, CA1300, and CA1500.

[0125] Table 8. Physical parameters for XRD peak fitting of different hard carbons.

[0126]

[0127] Figure 13 Table 8 shows the quantitative phase distribution obtained by XRD peak fitting. Specifically, the proportion of crystalline graphite phase in CA1100 is the highest (approximately 27.9%), which is consistent with the highly ordered layered structure observed in high-resolution transmission electron microscopy (HRTEM), indicating a significant tendency towards graphitization. As the temperature increases to 1300... o C and 1500 o C. The overall phase composition gradually shifts towards a pseudo-graphite vortex structure. The crystalline graphite phase in CA1500 decreases slightly (approximately 20.9%), while the pseudo-graphite phase increases to 75%. Although the proportion of highly disordered phases increases, they still constitute a minority, indicating that under high-temperature conditions, some originally ordered graphite regions evolve into pseudo-graphite structures with vortex packing. This leads to a unique phenomenon: "an increase in locally ordered regions, but the packing sequence becomes looser," rather than simple "complete disorder." CA1300 is mainly composed of pseudo-graphite structures (accounting for over 90%), containing only a small amount of highly disordered phases. This structural state balances, to some extent, moderate disorder and a continuous conductive network, which is more advantageous for achieving high cycling stability and high specific capacity in sodium-ion storage processes.

[0128] Figure 14 These are the fitted Raman spectra of CA1100, CA1300, and CA1500.

[0129] Table 9. Physical parameters of Raman fitting for different hard carbons.

[0130]

[0131] Figure 14 The data in Table 9 further confirms the difference in defect density. Specifically, the I of CA1300... D1 / I G and I D3 / IG The ratios were 1.53 and 0.62, significantly higher than those of CA1100 (1.38 and 0.26) and CA1500 (1.39 and 0.20). This indicates that CA1300 has a higher defect density and stronger short-range disorder, providing more active adsorption and storage sites for sodium ions and accelerating their diffusion kinetics. These results are in high agreement with the XRD phase analysis results, highlighting the unique advantages of CA1300 in optimizing the "order-disorder balance" and defect synergy.

[0132] Figure 15 These are high-resolution C1sXPS spectra of CA1100, CA1300, and CA1500. Figure 15 The C1s XPS spectra of the samples are presented, revealing differences in the composition of surface carbon species. All three spectra show a main peak at 284.5 eV, corresponding to sp2C-C / C=C bonds, but this peak varies with calcination temperature from 1100 eV. o C rises to 1500 o C, the oxygen-containing carbon peaks (CO at 286.0 eV and OC=O at 289.0 eV) gradually decrease. In the spectrum of CA1300, in addition to the main peak, the signals of CO and OC=O can still be clearly observed, indicating that a certain amount of oxygen-containing functional groups still exist on its surface. These functional groups help improve the wettability of the material with the electrolyte and promote the initial Na+ formation. + Adsorption and diffusion. In contrast, CA1100 has a higher content of CO and OC=O components, and the excessive oxygen sites promote the rapid formation of a thick SEI layer, consuming a large amount of Na. + This limits the rate of interfacial electrochemical reactions. Meanwhile, in CA1500, the oxygen-containing carbon peak is significantly weakened, leaving only a very weak OC=O component, indicating a highly deoxygenated and more inert surface. This may help suppress side reactions, but it also reduces polar sites and active adsorption centers that interact with the electrolyte.

[0133] Figure 16 Figure a shows the performance of CA1100, CA1300, and CA1500 electrodes at different current densities, and Figure b shows the charge-discharge curves of CA1300 at different voltages.

[0134] Figure 16This study compared the rate performance of anthracite-derived hard carbon at different calcination temperatures. As the current density increased from 0.1 A / g to 2.0 A / g, CA1300 maintained a significant reversible capacity, with capacities of 349.8, 266.7, 221.6, 196.1, 183.9, 182.6, and 153.8 mAh / g, respectively. This was significantly better than CA1100 and CA1500, demonstrating excellent rate performance. Notably, when the current density returned to 0.1 A / g, the capacity of CA1300 almost completely recovered to 348.5 mAh / g, with a recovery rate approaching 100%, indicating that at 1300... o The carbon framework and porous structure obtained under C2000 maintain excellent structural integrity and reversible sodium storage performance even after high-rate charge-discharge cycles, effectively promoting electron conduction and sodium diffusion. Furthermore, the galvanostatic charge-discharge (GCD) curves of CA1300 at different current densities exhibit a smooth slope characteristic without obvious plateau-related potential steps, indicating that the sodium storage mechanism is stable and consistent over a wide rate range. Figure 16 (b) Even at a high current density of 2.0 A / g, CA1300 still delivers a capacity of 155.6 mAh / g, with highly overlapping charge-discharge curves, confirming its excellent rate adaptability and structural stability. Furthermore, CA1300 exhibits superior specific capacity at various current densities compared to other hard carbon anode materials reported for sodium-ion batteries. For example, expanded graphite (AEG) from anthracite delivers a capacity of 326 mAh / g at 0.03 A / g and 140 mAh / g at 2.0 A / g, with a capacity retention of 82%; BC-LO-1300 has a specific capacity of 310 mAh / g at 0.1C and 70 mAh / g at 5C, with a capacity retention of 77.4%.

[0135] Figure 17 The cycling performance of CA1100, CA1300, and CA1500 electrodes was tested at a current of 0.1 A / g; and the charge-discharge curves of the CA1300 electrode at different cycle numbers at this current were also tested. Figure a shows the cycling performance of the CA1100, CA1300, and CA1500 electrodes at a current of 0.1 A / g; Figure b shows the charge-discharge curves of the CA1300 electrode at different cycle numbers at this current.

[0136] Figure 17This paper presents the cycling performance of the three samples at a current density of 0.1 A / g. After 250 cycles, CA1300 maintained a stable capacity of approximately 348.5 mAh / g, significantly higher than CA1100 (248.6 mAh / g) and CA1500 (289.3 mAh / g). In the initial cycles, all samples experienced capacity decline due to irreversible SEI film formation, active site deactivation, and microstructural changes in the carbon framework. Subsequently, the capacity tended to stabilize, with CA1300 exhibiting the best stability, indicating excellent structural stability and electrochemical reversibility during sodium ion intercalation / deintercalation. Figure 17 b shows that the GCD curve of CA1300 remained largely consistent from cycle 1 to cycle 100, with only a slight decrease in capacity. The capacity in the low-potential plateau region remained at approximately 194.2 mAh / g, reflecting the stable contribution of the low-potential plateau capacity even during long-term cycling.

[0137] Figure 18 These are the initial coulombic efficiency values ​​of the CA1100, CA1300, and CA1500 electrodes at a current of 0.1 A / g. Figure 18 The initial coulombic efficiency of the three electrodes at a current density of 0.1 A / g was compared. The results showed that the initial coulombic efficiency of CA1300 was 85.1%, significantly higher than that of CA1100 (73.9%) and CA1500 (78.6%). This indicates that at a current density of 1300... o The optimized carbon framework structure formed under C significantly improves the initial coulombic efficiency. This hierarchical porous structure and low surface oxygen content limit the interaction between the electrolyte and active sites, suppress side reactions, and promote the formation of a stable solid electrolyte interfacial film, thereby reducing capacity loss in the first cycle.

[0138] Figure 19 Figure 1 shows the first cycle charge-discharge curves of different samples at 0.1 A / g, and a comparison of the corresponding plateau capacity and ramp capacity. Specifically, Figure 2a shows the first cycle charge-discharge curves of different samples at 0.1 A / g; Figure 3b shows a comparison of the corresponding plateau capacity and ramp capacity.

[0139] Figure 19 Figure a shows the voltage plateaus of the three electrodes during the first charge-discharge cycle. Although the voltage plateaus of all three materials are similar in height and shape, the CA1300 electrode exhibits a higher reversible capacity ratio, indicating that its interfacial reaction has been optimized. Figure 19b shows the contribution ratio of the plateau and ramp regions in all samples. The CA1100 electrode exhibits the fastest reaction rate, which is mainly attributed to the role played by its dominant ramp region. With increasing calcination temperature, the defect density in the carbon layer decreases, thereby improving the plateau capacity and enhancing the reversible Na in coal-based hard carbon. + Storage performance. In the CA1300 electrode, the near 1:1 contribution ratio of the plateau and ramp regions balances high capacity and excellent rate performance.

[0140] Figure 20 The capacity retention rate of CA1100, CA1300 and CA1500 electrodes after 250 cycles at a current of 0.1 A / g. Figure 20 The capacity retention of these three electrodes after 250 cycles at a current density of 0.1 A / g is shown. The CA1300 electrode exhibits a capacity retention of 93.4%, significantly higher than the CA1100 electrode (90.1%) and the CA1500 electrode (91.3%). This indicates that at a current density of 1300... o The microstructure formed under C conditions, especially the high proportion of pseudo-graphite structure and moderate defect density, helps to maintain a stable framework and high electrochemical activity during long-term cycling.

[0141] Figure 21 This section describes the cycling performance of the CA1100, CA1300, and CA1500 electrodes at 1.0 A / g. Comparing the specific capacity and cycling performance of the three electrodes at a high current density of 1.0 A / g, it was found that CA1300 retained a capacity of 186.5 mAh / g after 1000 cycles, with a capacity retention of approximately 90.6%. Figure 21 This strongly demonstrates that its robust structural toughness and interfacial stability during sodium ion insertion / extraction can effectively suppress side reactions and degradation during long-term cycling.

[0142] Figure 22 Figure 1 shows the in-situ EIS spectra of the CA1100, CA1300, and CA1500 electrodes during their first cycle, along with the corresponding analog circuit diagrams. Specifically, Figure a shows the in-situ EIS spectra of the CA1100, CA1300, and CA1500 electrodes during their first cycle, while Figure d shows the corresponding analog circuit diagram.

[0143] Figure 22 The image shows the in-situ electrochemical impedance spectroscopy (EIS) of the CA1100, CA1300, and CA1500 electrodes during the first cycle. Based on the equivalent circuit diagram (… Figure 22 d), the semicircles observed at high and mid frequencies correspond to those attributed to the solid electrolyte interface (R). SEI The resistance of ) and the resistance attributed to charge transfer resistance (R) ctThe resistance of an electrolyte. The resistance caused by an electrolyte is represented by the symbol R. S express.

[0144] Figure 23 This is the AC impedance fitting resistance of the CA1300 electrode. Among the three electrodes, the CA1300 electrode has the lowest overall impedance, indicating that its ion / electron transport coupling is the most favorable, and its charge flow resistance within its microstructure is the smallest. Figure 23 This advantage is attributed to the optimized carbon framework formed at 1300 degrees Celsius, where moderate disorder and high-density structural defects synergistically improve interfacial charge transfer dynamics.

[0145] Figure 24 It is the AC impedance fitting resistor for the CA1100 electrode. Figure 25 This is the AC impedance fitting resistor for the CA1100 electrode. In contrast, the CA1100 electrode exhibits relatively high impedance across all frequencies and maintains a high R0 even at low potentials. ct and solid electrolyte interface resistance (R SEI This indicates that it is severely limited by interfacial and charge transfer resistance. This is due to its overly ordered structure, extremely few defects, and singular surface function, which restricts electron conduction and sodium ion transport. Figure 24 The resistance of CA1300 and CA1500 electrodes decreases as the potential decreases, with the CA1300 electrode showing the largest resistance decrease. Figure 25 This indicates that the CA1300 electrode exhibits the lowest kinetic barrier during deep sodiumization, enabling more efficient energy storage at low potentials. Overall, the CA1300 electrode demonstrates optimal impedance performance and the lowest interfacial and charge transfer resistance over the low potential range, explaining its excellent electrochemical performance and cycling stability.

[0146] Figure 26 These are in-situ EIS spectra of CA1100, CA1300, and CA1500 electrodes during their first cycle; RL of the three electrodes at different voltages. SEI Atlas. Figure 26 R of three electrodes at different potentials was compared. SEI The CA1300 electrode consistently exhibits a lower interface resistance than the CA1100 and CA1500 electrodes across the entire potential range, with a particularly significant advantage at low potentials. This reduces R... SEI It is beneficial for cycle stability because it reduces energy loss during charging and discharging and promotes faster ion / electron transport, thereby slowing performance degradation and improving long-term stability.

[0147] To further evaluate the diffusion behavior of sodium ions, Figure 27The GITT curves of the three electrodes are shown. Figure 27 These are the GITT curves for the CA1100, CA1300, and CA1500 electrodes. By observing the stability of the quasi-equilibrium plateau and the degree of polarization, the CA1300 electrode exhibits the smoothest plateau evolution and the least polarization, indicating the lowest diffusion resistance for sodium ion insertion / extraction and superior transport kinetics. In contrast, the CA1100 electrode shows a steeper and more volatile voltage response with the shortest plateau duration, indicating significant polarization and a slow diffusion process, which can be attributed to its highly ordered structure and limited defect-assisted channels for sodium ion transport. The performance of the CA1500 electrode falls between that of the CA1100 and CA1300 electrodes, but is still inferior to that of the CA1300 electrode.

[0148] Figure 28 Voltage-dependent sodium ion diffusion coefficient (logD) Na+ Quantitative comparisons were conducted. Figure 28 Na obtained through GITT + Diffusion coefficient (logD). The CA1300 electrode exhibits a significantly higher diffusion coefficient than the other two samples across most of the voltage range. Notably, the CA1300 electrode still shows a significant advantage in the low voltage region close to 0.01V, indicating that its structural features enable rapid sodium ion transport during deep sodification, thereby improving its practical sodium ion storage capacity. This rapid diffusion stems from the synergistic effect of multiple microstructural factors in the CA1300 electrode, including moderate interlayer spacing, a balanced long-range ordered / short-range disordered framework, abundant structural defects, and sufficient oxygen-containing surface functional groups, all of which contribute to sodium ion migration. In contrast, sodium ion diffusion in the CA1100 electrode is limited by its overly ordered framework structure and insufficient defect density. Although the CA1500 electrode benefits to some extent from increased disorder, excessive structural relaxation does not create the optimal permeation diffusion path, and its diffusion performance remains lower than that of the CA1300 electrode. In summary, the EIS and GITT results consistently demonstrate that the CA1300 electrode offers the most favorable kinetics and transport characteristics, highlighting its great potential as an anode material for sodium-ion batteries.

[0149] Figure 29 These are the CV curves of CA1100, CA1300, and CA1500 samples at scan rates of 0.1–1.0 mV / s. Figure a shows the CV curves of CA1100, Figure b shows the CV curves of CA1300, and Figure c shows the CV curves of CA1500 samples at scan rates of 0.1–1.0 mV / s.

[0150] Figure 29The results showed that when the scan rate increased from 0.1 mV / s to 1.0 mV / s, all samples exhibited distinct anodic and cathodic peaks, corresponding to the sodium ion insertion / extraction and extraction processes, respectively. With increasing scan rate, the peak current of all electrodes increased significantly, while the peak potential shifted slightly, indicating that the sodium ion insertion / extraction kinetics are related to the scan rate, and the electrochemical response becomes more rapid at higher scan rates. Notably, the CA1300 electrode exhibited the highest peak current density, indicating that its electrochemical activity is superior to the other two electrodes. Figure 29 (b) Furthermore, the reversibility of the reaction is also stronger. These results further confirm the advantages of the CA1300 microstructure in enhancing electrochemical performance.

[0151] To quantitatively distinguish the contributions of capacitive control and diffusion control to the total current response, the relationship i(v) = k1ν + k2ν was used. 1 / 2 To evaluate the relative proportions of capacitive and diffusion processes at different scan rates. Figure 30 This is a graph showing the proportional contribution of capacitance to diffusion control for CA1100, CA1300, and CA1500 at different scan rates. Figures a, b, and c are respectively graphs showing the proportional contribution of capacitance to diffusion control for CA1100, CA1300, and CA1500 at different scan rates.

[0152] Figure 30 The capacitance and diffusion effects of the three samples were compared at different scan rates. As the scan rate increased from 0.1 mV / s to 1.0 mV / s, the capacitance effect of all three materials increased significantly. The CA1300 electrode showed the largest increase, from 28.87% to 56.44%, indicating a stronger capacitance-dominated charge storage characteristic, resulting in superior performance under rapid charge-discharge conditions. The capacitance effect of the CA1100 electrode also increased from 25.73% to 51.15%, although its increase was smaller than that of the CA1300 electrode. In contrast, the CA1500 electrode consistently exhibited the lowest capacitance effect, increasing only from 13.25% to 32.14%, meaning that capacitance storage played the smallest role in this electrode.

[0153] Subsequently, the invention was further verified using cyclic voltammetry (CV) curves (scan rate of 0.6 mV / s). Figure 31 Figures a, b, and c show the capacitive contribution analysis of CA1100, CA1300, and CA1500 at a scan rate of 0.6 mV / s.

[0154] Figure 31The figures ac show the capacitance contribution of each electrode at this scan rate, with the shaded area representing the current controlled by capacitance. For the CA1100 and CA1300 electrodes, the calculated capacitance contributions are 46.65% and 48.81%, respectively, while the value for the CA1500 electrode is significantly lower at 24.96%, a marked decrease compared to the other two electrodes. This observation confirms that at a moderate calcination temperature (1100... o C and 1300 o Hard carbon electrodes prepared under C) have a higher capacitance contribution, making them more suitable for fast charge / discharge processes, while high-temperature electrodes (1500) o C) Storage processes are primarily diffusion-controlled, resulting in relatively slow energy storage kinetics. This difference stems from microstructural changes induced by calcination temperature. Higher temperatures (e.g., 1500 °C) lead to slower energy storage. o C) Increased graphitization reduces capacitance contribution. Lower structural disorder and fewer defect- and pore-related active sites hinder ion transport and limit surface-driven storage capacity, thus weakening capacitance. Appropriate moderate temperatures (e.g., 1300 °C) further contribute to capacitance. o C) It can balance the microstructure, improve electrochemical kinetics, and enhance capacitance, which explains its excellent fast charge and discharge performance.

[0155] Commercially available sodium vanadium phosphate (Na3V) was selected. 2( Na₃V₂(PO₄)₃ was used as the cathode material. Before assembling the full cell, the CA1300 anode was pre-charged using sodium metal as the counter electrode. The pre-charging process involved 10 charge-discharge cycles at a current density of 0.1 A / g. After pretreatment, the half-cell was carefully disassembled in a glove box, and the pre-charged CA1300 anode was paired with the Na₃V₂(PO₄)₃ cathode to assemble a complete sodium-ion full cell. The capacity ratio of the cathode to the anode was controlled at approximately 1:1.2. Battery performance was tested using a CT3002A battery testing system. Constant current charge-discharge tests were performed at different current densities within the voltage range of 0.5–3.8 V. In addition, cyclic voltammetry (CV) curves within the voltage range of 0.5–3.8 V were recorded using a CHI760E electrochemical workstation.

[0156] Figure 32 This is the CV curve of NVP / / CA1300 total sodium ions. Figure 32 The first three cyclic voltammetry curves of the full cell are shown, in which a pair of distinct anodic and cathode peaks are clearly observed. The anodic peak at 3.45 V corresponds to the Vo peak at the NVP cathode. 3+ Transform into V 4+ This process involves sodium ion extraction, while the cathode peak at 3.30V coincides with V. 4+The reduction and sodium ion insertion processes are correlated. The cyclic voltammetry curves of these three consecutive scans almost overlap, and the anodic / cathode peaks are highly symmetrical, indicating excellent reversibility and strong electrochemical stability even in the initial cycle. This behavior not only reflects a highly reversible redox reaction but also confirms the structural stability of the electrode material during repeated electrochemical processes.

[0157] In the rate performance test, Figure 33 The reversible capacity of the NVP / / CA1300 full cell at different current densities was demonstrated. Figure 33 This is a graph showing the rate performance test results of the NVP / / CA1300 all-sodium-ion battery at different current densities. As the current density increases from 0.05 A / g to 2.0 A / g, the reversible capacity gradually decreases from 186.6 mAh / g to 68.2 mAh / g. When the current density decreases again to 0.1 A / g, the capacity rapidly recovers to its initial level, demonstrating not only excellent rate performance but also strong capacity recovery capability. This result indicates that the CA1300 hard carbon anode can effectively maintain good full-cell performance under high-rate conditions, especially excelling in reversible capacity recovery.

[0158] Figure 34 These are the charge-discharge curves of the NVP / / CA1300 all-sodium ion battery at different current densities. Figure 34 The charge / discharge curves at different current densities are shown. At low current densities, the battery exhibits a relatively long voltage plateau region; as the current density increases, the plateau becomes shorter and steeper, indicating that sodium storage behavior is influenced by both capacitance-controlled and diffusion-controlled processes. Importantly, no significant voltage polarization was observed, and the capacity output remained stable throughout the charge / discharge process, demonstrating good kinetic performance and reliable energy storage behavior at high current densities, which is crucial for fast charge / discharge applications.

[0159] Figure 35 These are the cycle stability and coulombic efficiency at 1.0 A / g for 900 cycles. Figure 35 The results further demonstrate the stability and coulombic efficiency of the full cell during long-term cycling at a current density of 1.0 A / g. Although a slight capacity decay occurs in the initial few cycles, the capacity quickly stabilizes, maintaining a capacity of 84.6 mAh / g after 900 cycles. Simultaneously, the coulombic efficiency rises rapidly in the initial stage and remains close to 100% thereafter, indicating a high degree of reversibility of the electrochemical reactions and virtually no side reactions during long-term cycling. More importantly, this performance implies excellent interfacial stability within the battery, effectively suppressing side reactions and maintaining overall electrochemical integrity during long-term operation.

[0160] This invention also investigated the influence of a single key factor. During the research, based on the scheme of Example 1, the origin of the preferred conditions in steps 1, 2, and 3 of Example 1 was studied.

[0161] The important single-factor experiment in step 1 of Example 1: This experiment mainly studies the effects of HNO3 solution concentration, solid-liquid ratio of anthracite powder to HNO3 solution, and soaking time in HNO3 solution.

[0162] The effect of HNO3 solution concentration on the experiment is as follows: In step 1 of Example 1, HNO3 solutions of 1 mol / L, 2 mol / L, and 3 mol / L were used for soaking, and the soaking time was 6 h for each. Next, steps 2 and 3 of Example 1 were performed to prepare anthracite-based hard carbon. The obtained samples were named CA1300 / 1MHNO3, CA1300 / 2MHNO3, and CA1300 / 3MHNO3 according to the nitric acid concentration, respectively.

[0163] Table X1 below shows the NMR spectral results of the intermediate products after treatment with three concentrations of nitric acid and execution of step 2. This invention specifically detected the content of three oxygen-containing groups—hydroxyl, carboxyl, and carbonyl—in the intermediate products; the unit for each oxygen-containing group content is (mmol / g). The intermediate products treated with the three concentrations of nitric acid were named anthracite 1MHNO3, anthracite 2MHNO3, and anthracite 3MHNO3, respectively.

[0164] Table X1 summarizes the content of hydroxyl, carboxyl, and carbonyl groups in the intermediate products of anthracite 1MHNO3, anthracite 2MHNO3, and anthracite 3MHNO3.

[0165]

[0166] Table X2 below shows a summary of the ash content and silicon content test results in the intermediate product after treatment with three concentrations of nitric acid and execution of step 2.

[0167] Table X2 is a summary table of the test results for ash content and silicon content of intermediate products of anthracite 1MHNO3, anthracite 2MHNO3, and anthracite 3MHNO3.

[0168]

[0169] The analysis is as follows: Under the same solid-liquid ratio and soaking time, the content of oxygen-containing groups in the intermediate product of anthracite powder after two-step acid treatment gradually increases with the increase of nitric acid concentration. Furthermore, the ash content in this intermediate product gradually decreases with the increase of nitric acid concentration; however, under the same hydrofluoric acid conditions, the removal capacity of silica is comparable.

[0170] This experiment calculated the Ig of three hard carbon materials: CA1300 / 1M HNO3, CA1300 / 2M HNO3, and CA1300 / 3M HNO3. D / I G The values ​​are then summarized in Table X3 below.

[0171] Table X3, CA1300 / 1M HNO3, CA1300 / 2MHNO3 and CA1300 / 3MHNO3 are three types of hard carbon materials. D / I G Value summary table.

[0172]

[0173] The analysis is as follows: Under the same carbonization conditions, as the content of oxygen-containing functional groups such as hydroxyl, carboxyl, and carbonyl groups in the anthracite powder intermediate product increases after two-step acid treatment, the degree of anthracite-based graphitization decreases. D / I G The value increases. This is because the introduction of oxygen-containing functional groups such as hydroxyl and carboxyl groups creates steric hindrance between aromatic rings, inhibiting the directional stacking of aromatic units during carbonization, thereby reducing the degree of graphitization. In addition, as the ash content in this intermediate product decreases, its ability to catalyze graphitization weakens, resulting in a decrease in the degree of graphitization of anthracite-based coal.

[0174] This experiment tested the electrochemical performance of three hard carbon anodes: CA1300 / 1MHNO3, CA1300 / 2MHNO3, and CA1300 / 3MHNO3, and the results are summarized in Table X4 below.

[0175] Table X4 summarizes the electrochemical performance tests of three types of hard carbon anodes: CA1300 / 1MHNO3, CA1300 / 2MHNO3, and CA1300 / 3MHNO3.

[0176]

[0177] The analysis is as follows: Comparison reveals that the initial discharge specific capacity of the three types of hard carbon increases with decreasing graphitization. This is because hard carbon with low graphitization has a higher content of heteroatoms such as oxygen atoms, a higher defect content, and an increased number of sodium-storing active sites, thus increasing the initial sodium intercalation capacity. Furthermore, the initial coulombic efficiency of the three types of hard carbon exhibits a trend of first increasing and then decreasing with decreasing graphitization. This is because while increasing graphitization is beneficial for increasing the carbon interlayer spacing, thereby reducing sodium-storing dead zones and improving initial coulombic efficiency, excessively low graphitization leads to increased heteroatom and defect content, an increase in irreversible sodium-storing sites, an increase in irreversible capacity, and a decrease in initial coulombic efficiency.

[0178] Based on the above experimental results, a nitric acid concentration of 2 mol / L is the preferred condition. Next, with the nitric acid concentration fixed at 2 mol / L, the solid-liquid ratio of anthracite powder to HNO3 solution in step 1, as well as the soaking time in the HNO3 solution, will be optimized.

[0179] The solid-liquid ratio experiment of anthracite powder and HNO3 solution was conducted as follows: The solid-liquid ratio of anthracite powder to HNO3 solution in Step 1 was changed to 1:4, 1:6, and 1:8, respectively. After performing Step 2, the content of hydroxyl, carboxyl, and carbonyl groups in the intermediate products was detected. The intermediate products obtained with the three solid-liquid ratios were named anthracite 1:4 HNO3, anthracite 1:6 HNO3, and anthracite 1:8 HNO3, respectively. The resulting samples were named CA1300 / 1:4 HN, CA1300 / 1:6 HN, and CA1300 / 1:8 HN, respectively.

[0180] The HNO3 solution soaking time experiment was conducted as follows: The HNO3 solution soaking time in Step 1 was changed to 4h, 6h, and 8h respectively; after performing Step 2, the content of hydroxyl, carboxyl, and carbonyl groups in the intermediate products was detected; the intermediate products obtained at the three soaking times were named anthracite 4hHNO3, anthracite 6hHNO3, and anthracite 8hHNO3, respectively. The obtained samples were named CA1300 / 4hHN, CA1300 / 6hHN, and CA1300 / 8hHN, respectively.

[0181] Table X5 below shows the oxygen-containing group content (mmol / g), ash content, and silicon content in the intermediate products under different solid-liquid ratios and soaking times.

[0182] Table X5 is a summary table of the detection of oxygen-containing group content, ash content and silicon content in intermediate products under different solid-liquid ratios and soaking times.

[0183]

[0184] The analysis is as follows: With the increase of the solid-liquid ratio of anthracite powder and nitric acid and the extension of nitric acid soaking time, the content of oxygen-containing groups in the intermediate product of anthracite powder after two-step acid treatment gradually increases. Furthermore, with the increase of nitric acid concentration, the ash content in this intermediate product gradually decreases, but under the same hydrofluoric acid conditions, the removal capacity of silicon dioxide is comparable.

[0185] Table X6 shows the I content of anthracite-based hard carbon prepared under different solid-liquid ratios and soaking times. D / I G Values ​​and electrochemical performance.

[0186] Table X6. I-values ​​of anthracite-based hard carbon prepared under different solid-liquid ratios and soaking times. D / I GSummary table of values ​​and electrochemical performance.

[0187]

[0188] The analysis is as follows: Similarly, as the solid-liquid ratio of anthracite powder and nitric acid increases and the nitric acid soaking time lengthens, the content of oxygen-containing functional groups gradually increases, while the ash content gradually decreases, corresponding to the anthracite-based hard carbon I... D / I G As the value increases, the graphitization degree of anthracite-based hard carbon decreases. Therefore, with the increase of the solid-liquid ratio of anthracite powder and nitric acid and the increase of nitric acid soaking time, the initial sodium intercalation capacity of anthracite-based hard carbon increases, and the initial coulombic efficiency also shows a trend of first increasing and then decreasing.

[0189] Based on the above experimental results, the optimal conditions in step 1 are: a fixed nitric acid concentration of 2 mol / L, a solid-liquid ratio of anthracite powder to HNO3 solution of 1:6, and a soaking time of 6 h. Next, the hydrofluoric acid treatment conditions in step 2 will be optimized, specifically investigating the effects of different concentrations of hydrofluoric acid solution, different soaking temperatures, different mass-to-volume ratios of metal-removing anthracite powder to hydrofluoric acid solution, and different hydrofluoric acid soaking times.

[0190] It should be noted that, as demonstrated by the preceding experiments, hydrofluoric acid has a very weak effect on the oxygen-containing group content of anthracite in this invention, and mainly affects the silicon content and ash content in the material.

[0191] In experiments using hydrofluoric acid solutions of different concentrations, 1 mol / L, 2 mol / L, and 3 mol / L hydrofluoric acid solutions were used in step 2 of Example 1 to compare the effects of different concentrations of hydrofluoric acid treatment. The ash content and silicon content of the intermediate products anthracite 1HF, anthracite 2HF, and anthracite 3HF were also measured. The final products were named CA1300 / 1HF, CA13002HF, and CA1300 / 3HF, respectively.

[0192] Subsequently, in experiments at different immersion temperatures, with the hydrofluoric acid solution concentration fixed at 2 mol / L, immersion temperatures of 30℃, 50℃, and 70℃ were used to compare the effects of different immersion temperatures under the same concentration of hydrofluoric acid. The ash content and silicon content of the intermediate products, anthracite HF at 30℃, anthracite HF at 50℃, and anthracite HF at 70℃, were also measured. The final products were named CA1300 / 30℃HF, CA1300 / 50℃HF, and CA1300 / 70℃HF, respectively.

[0193] Subsequently, with the hydrofluoric acid solution concentration fixed at 2 mol / L and the soaking temperature fixed at 50℃, anthracite powder (excluding metal) was prepared with hydrofluoric acid solution at mass-to-volume ratios of 1 g:4 mL, 1 g:6 mL, and 1 g:8 mL, respectively. The effect of different ratios of anthracite powder to hydrofluoric acid solution under the same hydrofluoric acid concentration and soaking temperature was compared. The ash content and silicon content in the intermediate products anthracite 1:4HF, anthracite 1:6HF, and anthracite 1:8HF were also measured. The final products were named CA1300 / 1:4HF, CA1300 / 1:6HF, and CA1300 / 1:8HF, respectively.

[0194] Finally, with the hydrofluoric acid solution concentration fixed at 2 mol / L, the soaking temperature fixed at 50℃, and the anthracite powder (excluding metal) and hydrofluoric acid solution prepared at a mass-to-volume ratio of 1 g:6 mL, the effects of different soaking times were compared and studied. Specifically, the soaking times compared were 4 h, 6 h, and 8 h. The ash content and silicon content in the intermediate products anthracite 4hHF, anthracite 6hHF, and anthracite 8hHF were determined. The final products were named CA1300 / 4hHF, CA1300 / 6hHF, and CA1300 / 8hHF, respectively.

[0195] Table X7. Summary of ash and silicon content in intermediate products under different hydrofluoric acid treatment conditions.

[0196]

[0197] The analysis is as follows: A comparison revealed that during the hydrofluoric acid treatment process, as the hydrofluoric acid concentration, treatment temperature, solid-liquid ratio, and treatment time increased, the ash content in the intermediate product of the anthracite powder after the first two acid treatments first decreased and then stabilized. The silicon content also showed the same trend. Considering cost, the optimal hydrofluoric acid concentration was determined to be 2 mol / L, the treatment temperature to be 50℃, the solid-liquid ratio to be 1 g:6 mL, and the treatment time to be 6 h.

[0198] Table X8. I of anthracite-based hard carbon prepared under different hydrofluoric acid treatment conditions D / I G Summary table of values ​​and electrochemical performance.

[0199]

[0200] Analysis is as follows: Comparison reveals that, under the same oxygen-containing group content, the degree of graphitization of the material and the ash content in the intermediate product after the first two acid treatments of anthracite powder exhibit similar power changes. With increasing hydrofluoric acid concentration, treatment temperature, solid-liquid ratio, and treatment time, both show a trend of first decreasing and then stabilizing. Furthermore, the initial discharge capacitance of the material also shows a trend of first increasing and then slightly decreasing. However, the initial coulombic efficiency increases with increasing silicon content in the sample. This is because Na+ is generated during the initial sodium intercalation process. x SiO x And Na2O, which reduces the initial coulombic efficiency.

[0201] Based on the above experimental results, the optimal conditions for step 2 are: a hydrofluoric acid concentration of 2 mol / L, a soaking temperature of 50℃, a mixture of metal-free anthracite powder and hydrofluoric acid solution at a mass-to-volume ratio of 1 g:6 mL, and a soaking time of 6 h. Next, the effects of carbonization temperature and carbonization time in step 3 will be optimized.

[0202] Carbonization temperature experiment: The carbonization temperature in step 3 of Example 1 was changed to 1100℃, 1300℃, and 1500℃ respectively; the final product I was studied. D / I G Values ​​and electrochemical performance. The final products were named CA1100 (carbonized), CA1300 (carbonized), and CA1500 (carbonized), respectively.

[0203] Carbonization time experiment: The carbonization temperature was fixed at 1300℃; the carbonization time in step 3 was changed to 0.5h, 2h, and 4h; the final product's I... D / I G Values ​​and electrochemical performance. The final products were named CA1300 / 0.5h (carbonized), CA1300 / 2h (carbonized), and CA1300 / 4h (carbonized), respectively.

[0204] Table X9. I of anthracite-based hard carbon prepared under different carbonization conditions D / I G Summary table of electrochemical performance.

[0205]

[0206] The analysis is as follows: Under the same oxygen-containing groups and ash conditions, the degree of graphitization increases with increasing carbonization temperature. D / I G The initial discharge capacity decreases as the initial discharge rate decreases. The initial coulombic efficiency, however, shows a trend of first increasing and then decreasing. Ultimately, the optimal carbonization temperature was determined to be 1300℃ and the optimal carbonization time to be 2 hours.

[0207] It is understood that the present invention has been described through some embodiments, and those skilled in the art will recognize that various changes or equivalent substitutions can be made to these features and embodiments without departing from the spirit and scope of the invention. Furthermore, under the teachings of the present invention, these features and embodiments can be modified to adapt to specific situations and materials without departing from the spirit and scope of the invention. Therefore, the present invention is not limited to the specific embodiments disclosed herein, and all embodiments falling within the scope of the claims of this application are within the protection scope of the present invention.

Claims

1. A method for preparing anthracite-based hard carbon sodium-ion battery anode material, characterized in that, include Step 1: Soak anthracite powder in HNO3 solution at room temperature to remove some metal elements and introduce some oxygen-containing functional groups, including hydroxyl, carboxyl and carbonyl groups; after the HNO3 solution soaking treatment is completed, filter out the solid filter residue, then wash the filter residue with distilled water until neutral and dry it to obtain metal-free anthracite powder. The concentration of the HNO3 solution used was 1~3 mol / L. The anthracite powder and HNO3 solution were prepared at a mass-volume ratio of 1g:4~8 mL, and the soaking time was 4~8 h. Step 2: Soak the anthracite powder obtained in Step 1 in hydrofluoric acid solution at a temperature of 30~70℃ to remove some of the silica; after the hydrofluoric acid solution soaking treatment is completed, filter out the filter residue, then wash the filter residue with distilled water until neutral and dry it to obtain silica-free anthracite powder. The concentration of the hydrofluoric acid solution used is 1~3 mol / L. Except for the metal anthracite powder, the hydrofluoric acid solution is prepared at a mass-volume ratio of 1g:4~8mL, and the soaking time is 4~8h. Step 3: Heat and carbonize the desiliconized anthracite powder in a tube furnace; under a nitrogen atmosphere, heat to 1100~1500 ℃ at a heating rate of 5℃ / min and hold at this temperature for 0.5~4 h; then cool naturally to room temperature to obtain anthracite-based hard carbon sodium-ion battery anode material.

2. The method for preparing the anthracite-based hard carbon sodium-ion battery anode material according to claim 1, characterized in that, Anthracite powder is made from powder with a mesh size between 200 and 250.

3. The method for preparing the anthracite-based hard carbon sodium-ion battery anode material according to claim 1, characterized in that, In step 1, the concentration of the HNO3 solution used is 2 mol / L, and the anthracite powder and HNO3 solution are prepared at a mass-volume ratio of 1g:6mL, and the soaking time is 6 h; In step 2, the anthracite powder obtained in step 1 was soaked in hydrofluoric acid solution at a temperature of 50°C. The concentration of the hydrofluoric acid solution used was 2 mol / L. The anthracite powder and hydrofluoric acid solution were prepared at a mass-volume ratio of 1 g: 6 mL, and the soaking time was 6 h. In step 3, the material is heated to 1300 °C at a heating rate of 5 °C / min under a nitrogen atmosphere and held at this temperature for 2 h; then it is naturally cooled to room temperature to obtain anthracite-based hard carbon sodium-ion battery anode material.

4. An application of an anthracite-based hard carbon sodium-ion battery anode material, characterized in that, The anthracite-based hard carbon sodium-ion battery anode material prepared by the method described in claim 1 is used as the anode material for sodium-ion batteries.

5. The application of the anthracite-based hard carbon sodium-ion battery anode material according to claim 4, characterized in that, The working electrode is fabricated as follows: The working electrode is prepared by mixing 80 wt.% hard carbon, 10 wt.% Super-P carbon black, and 10 wt.% sodium alginate in an aqueous solution; the electrode's mass loading is 1.8~2.6 mg / cm³. 2 Whatman glass fiber was used as the diaphragm, and 1M NaPF6 dissolved in ethylene glycol dimethyl ether = 100 vol% was used as the electrolyte. All components are vacuum dried for 12 hours to remove moisture; the gas in the glove box must be replaced more than three times before operation to ensure an inert atmosphere; the water and oxygen content in the glove box is <1ppm, the vacuum drying oven is 60-80℃, and the tablet press pressure is 50MPa; Assembly sequence: Start with positive electrode shell, positive electrode shell → positive electrode plate → electrolyte → separator → sodium plate → gasket → spring plate → negative electrode shell. The electrolyte injection error must be controlled within ±0.02mL, and the separator must be completely soaked. Use a digital display tablet press to seal at 50MPa pressure to ensure contact resistance <0.5mΩ.

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