English

• 0.   Introduction

  Lithium‑ion batteries (LIBs) have been widely used in consumer electronics, electric vehicles, and energy storage devices due to their large capacity, good rate capability, and high cycling stability^[1-5]. The commercial graphite negative electrodes have limited the further increase in energy density of LIBs due to their relatively low theoretical capacity^[5-7]. Silicon is considered the mainstream negative electrode material for the next generation of LIBs due to its high theoretical specific capacity, low lithium removal/insertion potential platform, and abundant natural reserves^[6-8]. However, the low intrinsic electronic conductivity of silicon materials, as well as the structural damage caused by significant volume changes (beyond 300%) during lithium removal/insertion processes and the thickening of unstable solid electrolyte interface (SEI) films, limit its commercial application as negative electrode material for LIBs^[9-11].

  To solve the above-mentioned problems of silicon-negative electrodes, researchers have made tremendous efforts. Carbon coating is a common modification method for silicon-negative electrodes. The carbon coating not only increases the electronic conductivity of silicon-based materials and suppresses their volume expansion, but also avoids direct contact between the silicon substrate and electrolyte, and stabilizes the SEI film, thus increasing the electrochemical cycling stability of silicon-based electrodes^[12-15]. Introducing metal elements is another important method to improve the electrochemical lithium storage performance of silicon-based materials^[16-22]. Combining silicon materials with metal elements can improve the electronic conductivity of silicon-based materials^[19]. On the other hand, it is beneficial to alleviate the enormous stress of silicon during lithiation reactions, and has a positive effect on maintaining the electrochemical cycling stability of silicon anodes^[20]. The introduced metal elements can be divided into inactive lithium storage elements (e.g., Cu, Ag) and active lithium storage elements (e.g., Sn, Sb)^[19-21]. Elements with lithium storage activity can contribute to electrochemical capacity, and since different active materials generally have different lithium insertion/delithiation potentials, they can act as buffer layers to alleviate the volume expansion generated during their respective lithiation reactions^[20-22]. To avoid the pulverization and fragmentation of the silicon matrix caused by its huge volume expansion during the lithium insertion process, in addition to reducing the particle size of silicon to below the critical fracture size of 150 nm, the influence of volume expansion can also be buffered by constructing a porous structure^[23-26]. The rich pore structure inside porous silicon can not only shorten the transmission path of lithium ions but also provide sufficient buffer space for the volume change of silicon during cycling, thereby enhancing the electrochemical lithium storage performance of porous silicon-based materials^[26-29]. The strategies of carbon coating, introduction of metal elements, and construction of porous structures mentioned above have to some extent improved the low electronic conductivity and poor electrochemical cycling performance caused by significant volume change during lithium removal/insertion of silicon negative electrodes. However, how to further improve the electrochemical lithium storage performance of silicon-based materials is still a challenging task^[7]. At present, there are few literature reports on the combined application of carbon coating, introduction of metal elements, and construction of porous structures to improve the electrochemical lithium storage performance of silicon-based materials.

  We first used an alkaline etching and chemical displacement one-step method to deposit a Zn-Sb-Sn layer on the surface and inside of porous Si-Al microspheres (pSi-Al/Zn-Sb-Sn), with industrial grade SiAl alloy as raw material. Then, after heat treatment and acid etching, porous silicon/antimony-tin material (pSi/Sb-Sn) was obtained. Finally, the final product (pSi/Sb-Sn@C) was obtained by dopamine coating and high-temperature calcination treatment. The as-constructed silicon-based composite exhibited excellent characteristics of active metal components, carbon coating layer, and porous structure, demonstrating significantly enhanced electrochemical lithium storage performance.

  1.   Experimental

  1.1   Preparation of materials

  All chemicals except for SiAl alloy used are analytical grade and purchased from Sinopharm Chemical Reagent Co., Ltd. Bimetallic (Sb-Sn) surface-modified porous silicon microspheres (pSi/Sb-Sn) were prepared by alkaline etching and chemical displacement method. 1.0 g SiAl alloy powders (12% Si, ca. 325 mesh, provided by Alfa Aesar) were added to a 100 mL aqueous solution containing 1 mol·L^-1 KOH, 0.005 mol·L^-1 SbCl₃, 0.03 mol·L^-1 ZnCl₂, 0.06 mol·L^-1 ethylenediaminetetra-acetic acid (EDTA), and 0.002 5 mol·L^-1 SnCl₄·5H₂O, with continuous stirring. After the reaction for 5 min, pSi-Al/Zn-Sb-Sn precursor was collected by centrifugation, washed alternately with deionized water and ethanol three times, and dried at 70 ℃ for 10 h. The obtained precursor was annealed at 220 ℃ for 2 h in a tube furnace with an argon atmosphere. Then, after etching with 0.5 mol·L^-1 HCl, pSi/Sb-Sn was achieved via the same centrifugation, washing, and drying procedures.

  pSi/Sb-Sn@C composite was obtained by carbon coating treatment of pSi/Sb-Sn. 100 mg pSi/Sb-Sn was dispersed in 150 mL trihydroxymethylaminomethane hydrochloride (Tris) buffer solution (pH=8.5, 10 mmol·L^-1), and 60 mg dopamine hydrochloride was added under the condition of continuous stirring at room temperature. After the reaction proceeded for 20 h, the obtained reaction products were washed by centrifugation with deionized water and anhydrous ethanol and then dried under vacuum at 60 ℃ for 10 h. The final product pSi/Sb-Sn@C with a yield of ca. 98 mg was prepared by heating the as-obtained sample at 500 ℃ for 2 h in a tube furnace with an argon atmosphere.

  Using the same preparation method as pSi/Sb-Sn, porous silicon microspheres (pSi), and Sb-modified porous silicon microspheres (pSi/Sb) for control analysis were obtained by removing SbCl₃ and SnCl₄∙5H₂O or only removing SnCl₄∙5H₂O in the alkaline etching and chemical displacement solutions.

  1.2   Physical characterizations

  The microstructures and morphologies of the samples were characterized using field emission scanning electron microscopy (SEM, HITACHI SU8010, 5.0 kV, 10 μA) and high-resolution field emission transmission electron microscopy (HRTEM, JEM2100F, accelerated voltage 200 kV). The crystal structures of all samples were analyzed by X‑ray diffraction (XRD, Rigaku Ultima Ⅳ, 40 kV, 30 mA, Cu Kα radiation: λ=0.154 18 nm) at a scanning rate of 3 (°)·min^-1 within a scanning range of 20°-80°. The surface compositions and chemical states of the samples were characterized by X-ray photoelectron spectroscopy (XPS, ESCALAB 250Xi, 15 kV, 10 mA, Al Kα). Raman spectroscopy testing was conducted using a UV laser Raman spectrometer (LabRamHRUV) with a 532 nm laser emitter. Thermogravimetric analysis (TGA) was performed by a thermogravimetric analyzer (Netzsch STA409PC) in an air atmosphere, with a heating rate of 10 ℃·min^-1. The contents of Si, Sb, and Sn in the composite materials were determined using an inductively coupled plasma-optical emission spectrometer (ICP‑OES, OPTIMA 8000DV). The N₂ adsorption-desorption isotherms and pore size distributions of the samples were measured using a BET (Brunauer‑Emmett‑Teller) analyzer (Micromeritics ASAP 2020 M).

  1.3   Electrochemical measurements

  Active material, sodium alginate, and acetylene carbon black with a mass ratio of 70∶15∶15 were dispersed in deionized water to prepare a working electrode slurry with appropriate viscosity. The obtained slurry was coated on the copper foil collector and then dried at 60 ℃ for 12 h in a vacuum furnace to obtain the working electrode. The loading amount of silicon-based composite materials in the electrodes was ca. 1.0 mg·cm^-2. The assembly of the half cells (CR2025 buckle-type batteries) was carried out in a glove box filled with Ar, and the water/oxygen contents were strictly controlled below 10^-5. In the assembled CR2025 buckle‑type batteries, lithium foil was used as the counter electrode and reference one, and the as‑ prepared electrode and the Celgard 2325 polypropylene film were applied as the working electrode and separator, respectively. The electrolyte was prepared by dissolving 1 mol·L^-1 LiPF₆ in a mixed solution of methyl carbonate ethyl ester, dimethyl carbonate, and ethylene carbonate (1∶1∶1,V/V), adding fluorocarbonate ethylene (FEC, a mass fraction of 10%) as an additive. The assembled button-type half cells were aged at 45 ℃ for 24 h in a blast oven before electrochemical testing. Cyclic voltammetry (CV) curves were measured within the potential range of 0.01-3.00 V (vs Li⁺/Li) using a CHI 660D electrochemical workstation. Electrochemical impedance spectroscopy (EIS) was tested under open circuit potential with an amplitude of 10 mV and a frequency range of 10 mHz to 100 kHz, using an advanced electrochemical analyzer (Parstat 2273). Constant current charging-discharging tests were conducted using a NEWARE instrument within a potential window of 0.01-3.00 V at different current densities.

  2.   Results and discussion

  2.1   Preparation process and physical characterizations

  The preparation process of the pSi/Sb-Sn@C composite is illustrated in Fig. 1. Using inexpensive silicon aluminum alloy as raw material, pSi-Al/Zn-Sb-Sn intermediate was obtained through a one-step alkaline etching and chemical displacement reaction, involving the following chemical reactions:

  $ 2 \mathrm{Al}+6 \mathrm{H}_2 \mathrm{O}+2 \mathrm{OH}^{-} \rightarrow 2 \mathrm{Al}(\mathrm{OH})_4^{-}+3 \mathrm{H}_2 \uparrow $

  (1)

  $ \begin{aligned} \mathrm{Al}+\mathrm{M}-\mathrm{EDTA} \rightarrow \mathrm{Al}-\mathrm{EDTA}+\mathrm{M} \\ \\ \qquad(\mathrm{M}=\mathrm{Zn}, \mathrm{Sb}, \mathrm{Sn}) \end{aligned} $

  (2)

  Figure 1

  

  Figure 1.  Schematic illustration for the fabrication process of pSi and pSi/Sb-Sn@C composite

  In the naming of different materials, the symbol"/"represents the interface between the substrate and the metal deposition layer, and the symbol"-"is used to connect different components in the alloy or deposition layer.

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  The intermediate pSi-Al/Zn-Sb-Sn was subjected to heat treatment at 220 ℃ in an argon atmosphere and subsequent acid etching to obtain the precursor pSi/ Sb-Sn. During the heat treatment process, the diffusion of metal atoms and alloying reactions may occur, while Al and Zn with relatively high activity are preferentially dissolved during the acid etching. Using pSi/Sb-Sn as a precursor, the final product pSi/Sb-Sn@C composite was obtained through dopamine encapsulation and a high-temperature carbonization process. It should be noted that the introduction of zinc during the preparation process can avoid violent reactions caused by the large potential difference between Al and the substituted metals (Sb, Sn), thus facilitating the acquisition of ordered porous silicon microsphere structure^[30].

  The microstructure and surface element composition of the intermediate pSi-Al/Zn-Sb-Sn are shown in Fig. 2. As shown in Fig. 2a and 2b, the intermediate exhibited a completely porous microsphere structure. The element distribution mappings (Fig. 2c-2h) show that after the alkaline etching and chemical displacement process, metal Zn, Sb, and Sn were uniformly deposited inside and on the surface of the porous silicon sphere, while some metal aluminum components were also present.

  Figure 2

  

  Figure 2.  (a) TEM image, (b) HRTEM images, and (c-h) EDS elemental mappings of pSi-Al/Zn-Sb-Sn

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  Fig. 3 shows the SEM images of SiAl alloy, pSi/Sb-Sn, and pSi/Sb-Sn@C. The SiAl alloy precursor exhibited a spherical structure with a smooth surface and a diameter of 1-5 μm (Fig. 3a). The pSi/Sb-Sn particles retained the porous microsphere structure of the intermediate pSi-Al/Zn-Sb-Sn. Due to the removal of Al and Zn elements by the etching, the porous microspheres exhibited a more abundant pore structure (Fig. 3b-3d). It can be seen from Fig. 3d that the surface of pSi/Sb-Sn microspheres was relatively rough, which may be caused by the deposition of Sb/Sn metal particles on the surface of porous silicon microspheres. The pSi/ Sb-Sn@C sample obtained after carbon coating displayed a microsphere structure, and the surface pore structure disappeared (Fig. 3e, 3f).

  Figure 3

  

  Figure 3.  SEM images of SiAl alloy (a), pSi/Sb-Sn (b-d), and pSi/Sb-Sn@C (e, f)

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  The microstructure and element distribution information of the pSi/Sb-Sn@C sample were further characterized by TEM, and the results are presented in Fig. 4. As shown in Fig. 4a, pSi/Sb-Sn@C displayed regular porous microspheres and the interior of the microspheres presented a 3D network structure. HRTEM (Fig. 4b) showed clear lattice stripes, with a lattice spacing of 0.314, 0.311, and 0.307 nm attributed to the Si (111), Sb (012), and SbSn (101) crystal planes, respectively. It can also be observed in Fig. 4b that there is an uneven amorphous thin carbon layer on the outer surface area of the sample particles. From the surface distribution results of elements in Fig. 4c-4h, it can be seen that elements such as Sn, Sb, C, and N were relatively evenly distributed on the surface and inside of porous silicon microspheres, with N element originating from the precursor dopamine of carbon coating. It should be pointed out that microgrid-supported carbon film was used in TEM testing for sample preparation, resulting in a clear supporting carbon film skeleton shape in the C element distribution mapping. However, careful observation can identify the spherical carbon element distribution area of the sample.

  Figure 4

  

  Figure 4.  (a) TEM image, (b) HRTEM image, and (c-h) EDS elemental mappings of pSi/Sb-Sn@align="center"

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  To confirm the composition and microstructure of the final product (pSi/Sb-Sn@C), XRD, Raman, TGA, and N₂ adsorption-desorption tests were conducted, and the relevant results are shown in Fig. 5. Fig. 5a and 5b present the comparison between the XRD patterns of pSi and pSi/Sb-Sn@C. All diffraction peaks in the XRD patterns of pSi can be attributed to elemental Si (PDF No.27-1402). For pSi/Sb-Sn@C, besides the diffraction peaks of elemental Si and Sb (PDF No.35-0732), the weak diffraction peak at 29.1° (Fig. 5b) can be attributed to SbSn alloy (PDF No.33-0118), and no diffraction peaks of other substances were observed. This indicates that Sn and Sb were generated on the Si surface through the chemical displacement process, Reaction 2), and subsequent heat treatment leads to the formation of the SbSn binary alloy (Fig. 5b). No diffraction peaks of Sn and C can be found in the XRD patterns, possibly due to their poor crystallinity or lower content. The ICP-OES test results indicated that the atomic ratio of Si, Sb, and Sn metal components in pSi/Sb-Sn@C was 59.4∶37.4∶3.2.

  Figure 5

  

  Figure 5.  XRD patterns (a), enlarged XRD pattern (b), and Raman spectra (c) of pSi and pSi/Sb-Sn@C; TGA curves of pSi/Sb-Sn and pSi/Sb-Sn@C (d); N₂ adsorption-desorption isotherms of pSi (e) and pSi/Sb-Sn@C (f)

  Inset: the corresponding pore size distribution curves.

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  Fig. 5c shows the Raman spectra of the pSi and pSi/Sb-Sn@C samples. Both samples exhibited relatively strong peaks at 291, 510, and 930 cm^-1, which are attributed to elemental Si^[31-32]. The characteristic Raman peaks of the pSi/Sb-Sn@C composite at 1 344 and 1 588 cm^-1 are ascribed to the D and G peaks of carbon^[33-34], respectively. There are no D/G peaks attributed to carbon in the Raman spectrum of the pSi sample. This confirms the existence of carbon components with a certain degree of graphitization in the pSi/Sb-Sn@C sample. The TGA results of the pSi/Sb-Sn and pSi/Sb-Sn@C samples are shown in Fig. 5d. As the temperature increased, the Si, Sb, and Sn components in the two samples underwent oxidation to generate corresponding oxides, causing an increase in sample weight. For the pSi/Sb-Sn@C sample, the reaction of carbon oxidation to form CO₂ would simultaneously occur, which can lead to a decrease in the weight of the sample. Fig. 5d showed that when the temperature exceeded 300 ℃, the increase rate in sample weight of pSi/ Sb-Sn@C was significantly lower than that of the pSi/ Sb-Sn sample, which is the result of the combined effect of metal component oxidation and carbon gasification. By comparing the TGA curves of the two samples, the carbon mass fraction in the pSi/Sb-Sn@C sample can be calculated to be approximately 9.61%. Fig. 5e and 5f show the N₂ adsorption-desorption isotherms and pore size distribution mappings of the pSi and pSi/Sb-Sn@C samples, respectively. Two samples show type Ⅱ adsorption-desorption isotherms, indicating the presence of porous structure inside the two samples. The specific surface areas of the pSi and pSi/Sb-Sn@C samples were calculated to be 86 and 37 m²·g^-1 by the BET method, respectively. The pore size distribution mappings further indicate that pSi exhibited a wider pore size distribution and a larger pore volume ratio than pSi/Sb-Sn@C. The deposition of metal components in the pore structure and surface carbon coating leads to the narrow pore size distribution and relatively low specific surface area of the pSi/Sb-Sn@C sample.

  XPS technology was used to characterize the elemental composition and valence states of the pSi/ Sb-Sn@C sample, and the results are shown in Fig. 6. The XPS full spectrum (Fig. 6a) confirms the presence of Si, Sb, Sn, C, and O elements in the pSi/Sb-Sn@C sample, where O element is introduced due to surface oxidation of Si or other elements. In the high-resolution XPS spectrum of Si2p (Fig. 6b), the characteristic peaks at 102.4 and 103.2 eV correspond to Si—C and Si—O bonds^[35], respectively, while the characteristic peaks at 99.0 and 99.5 eV originate from Si—Si bond^[36]. The presence of Si—O bond indicates a certain degree of oxidation on the Si surface. The Si—C bond indicates that the substrate silicon has bonded with the coated carbon layer, enhancing their combination degree, which is beneficial for improving the electrochemical cycling stability of the composite. Fig. 6c shows the high-resolution XPS spectrum of C1s, where the strong characteristic peak at 284.8 eV corresponds to the C—C bond^[37], while the weak characteristic peaks at 286.0 and 289.0 eV are attributed to N—C and C=O bonds^[22,38], respectively. This indicates that C had a certain degree of oxidation, and the N element derived from the precursor has achieved doping with C, which is beneficial for improving the stability of Si/C composites^[22]. In the high-resolution XPS spectrum of Sn3d (Fig. 6d), the initial two peaks correspond to the Sn3d_3/2 and Sn3d_5/2 orbitals of Sn^[39], respectively. The Sn3d_3/2 peak of Sn can be divided into three peaks, with binding energies of 495.0 (Sn⁰), 495.8 (Sn²⁺), and 496.8 eV (Sn⁴⁺)^[40], respectively. The Sn3d_5/2 peak of Sn can also be decomposed into three peaks representing different valence states of Sn at 486.6 eV (Sn⁰), 487.0 eV (Sn²⁺), and 487.7 eV (Sn⁴⁺)^[40]. The results in Fig. 6d indicate that Sn had a certain degree of oxidation. Fig. 6e shows the high-resolution Sb3d XPS spectrum of pSi/Sb-Sn@C and pSi/Sb samples. The Sb3d_5/2 peak of the pSi/Sb sample can be decomposed into Sb⁰ (528.0 eV), Sb³⁺ (530.8 eV), and Sb—O (532.0 eV) peaks^[41], while its Sb3d_3/2 peak can be divided into Sb⁰ (537.4 eV) and Sb³⁺ (540.0 eV) peaks^[42]. Compared to the pSi/Sb sample, due to the introduction of Sn, the Sb³⁺ and Sb—O peaks of the pSi/Sb-Sn@C sample shifted by 0.5-0.7 eV towards the direction of high binding energy, while the peaks corresponding to Sb⁰ shifted by 2.5-2.7 eV towards the direction of high binding energy. The relatively larger shift of the Sb⁰ peak towards the direction of high binding energy may be due to the formation of the SbSn alloy phase^[43].

  Figure 6

  

  Figure 6.  XPS spectra of the pSi/Sb-Sn@C composite: (a) full survey, (b) Si2p, (c) C1s, and (d) Sn3d; (e) Sb3d XPS spectra of pSi/Sb and pSi/Sb-Sn@C composites

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  2.2   Electrochemical properties

  The electrochemical lithium storage properties of pSi/Sb-Sn@C and pSi samples were tested using electrochemical methods, and the relevant results are shown in Fig. 7. Fig. 7a shows the initial three CV curves of pSi/Sb‑Sn@C. A weak band centering at around 1.20 V appeared in the first cathodic scan, and it was absent in subsequent cycles, which may result from the irreversible reduction of fluoroethylene carbonate additive in the electrolyte^[15]. In the first cathode curve, the reduction peak at around 0.74 V corresponds to the irreversible formation of SEI film^[44] and the lithium insertion process of Sb^[45]. The reduction peak below 0.1 V corresponds to the process of lithium ions first embedding into crystalline silicon to form amorphous Li_xSi^[46]. In the first anodic curve, the oxidation peaks at 0.36 and 0.52 V correspond to the process of Li⁺ extraction from amorphous Li_xSi to amorphous silicon^[46], and the process of Li⁺ extraction from Li_ySb occurred at 1.07 V^[45]. In the second CV cycle curve, the peak potential of the cathode peak at 0.74 V on the first CV curve shifted positively to 0.81 V, while its peak current and peak area significantly decreased. This is due to the irreversible formation of the SEI film and the structure rearrangement of Sb in the first cycle^[44-46]. From Fig. 7a, it can be seen that as the number of CV cycles increased, the lithium insertion/ removal peak currents of Si gradually increased, corresponding to the electrochemical activation process of Si. Fig. 7b shows the initial three constant current charge-discharge curves of the pSi/Sb-Sn@C electrode at a current density of 100 mA·g^-1. During the initial lithium insertion process, there was a gentle potential plateau at approximately 0.1 V, corresponding to the formation of amorphous Li_xSi by lithium-ion insertion into crystalline silicon^[46]. This plateau disappeared during subsequent charging and discharging processes, which is consistent with the CV test results. The initial insertion and removal capacities of the pSi/Sb-Sn@C electrode were 2 745.6 and 2 062.3 mAh·g^-1, respectively, with a corresponding initial Coulombic efficiency of 75.1%. The initial capacity loss was mainly due to the irreversible formation of the SEI film. The Coulombic efficiency of the second charge-discharge cycle rapidly increased to 93.1%, and then stabilized at over 95%, demonstrating the stability of the SEI film.

  Figure 7

  

  Figure 7.  (a) Initial CV curves at a scan rate of 0.1 mV·s^-1 and (b) galvanostatic charge-discharge curves at 100 mA·g^-1 of pSi/Sb-Sn@C electrode; (c, d) Rate performance of pSi and pSi/Sb-Sn@C electrodes; (e) Cycle performance curves of pSi, pSi/Sb-Sn, and pSi/Sb-Sn@C electrodes

  The bar chart represents the discharge capacity, while the line diagram represents the retention rate vs 3rd cycle in c.

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  Fig. 7c and 7d show the rate capability of the pSi/Sb-Sn@C and pSi electrodes, where Fig. 7d is plotted in terms of the data in Fig. 7c, clearly presenting the average discharge capacity and retention rate of the two electrodes as a function of current density. The reversible capacities of the pSi/Sb-Sn@C electrode at current densities of 0.1, 0.2, 0.5, 1.0, 2.0, and 4.0 A·g^-1 were 2 249.4, 2 075.1, 1 817.2, 1 636.7, 1 408.9, and 1 054.5 mAh·g^-1, respectively. When the current density returned to 0.2 A·g^-1, the reversible capacity of the pSi/Sb‑Sn@C electrode was 1 842.5 mAh·g^-1, which is close to its initial capacity at the same rate. Compared to the pSi electrode, the pSi/Sb‑Sn@C electrode showed a significant increase in reversible capacities at all tested charge-discharge rates, indicating its excellent electrochemical lithium storage rate performance.

  Fig. 7e compares the electrochemical cycling performance of pSi, pSi/Sb-Sn, and pSi/Sb-Sn@C electrodes. Initial five charge-discharge activation cycles of the electrodes were performed at a small current of 100 mA·g^-1, and then the subsequent 295 electrochemical cycles were conducted at a relatively high current density of 1 A·g^-1. During the first 30 cycles at 1 A·g^-1, the capacity of the pSi/Sb-Sn@C electrode slowly increased, mainly due to the gradual activation of the electrode. After 300 charging-discharging cycles, pSi, pSi/Sb-Sn, and pSi/Sb-Sn@C electrodes delivered the reversible capacities of 418.4, 953.8, and 1 247.4 mAh·g^-1, respectively. In our previous work^[47], the carbon‑coated porous silicon microsphere composite prepared using a similar method to this work exhibited reversible capacities of 773.9 and 157.0 mAh·g^-1 after 200 charging-discharging cycles at 200 and 2 000 mA·g^-1 rates, respectively. The structural collapse of pSi induced by the huge volume change was responsible for its poor cyclic stability^[15,47]. The carbon coating or Sb-Sn surface deposition improved the capacity decay of pSi during electrochemical cycling to a certain extent. The synergistic effect of carbon coating and Sb‑Sn surface deposition significantly enhanced the electrochemical cycling performance of pSi.

  To investigate the electrochemical lithium storage kinetics of electrode materials, the EIS testing of the pSi and pSi/Sb-Sn@C electrodes after 300 charge‑ discharge cycles at a current density of 1 A·g^-1 was conducted, and the results are shown in Fig. 8a. The Nyquist plots of both electrodes were composed of a capacitance loop in the high-frequency region and an oblique line in the low-frequency region. The capacitance loops were caused by the relaxation process of charge transfer resistance (R_ct) and interface capacitance (CPE), while the diffusion process of lithium ions leads to the oblique line in the low-frequency region^[48]. The illustration inside Fig. 8a shows the corresponding equivalent circuit diagram. The parameter values obtained by fitting the Nyquist plots using the equivalent circuit diagram are listed in Table 1. The ohmic resistance (R_s), charge transfer resistance (R_ct), and diffusion impedance (Z_W) of the pSi/Sb-Sn@C electrode were significantly lower than the corresponding values of the pSi electrode. The surface deposition of Sb/Sn and carbon coating improved the electronic conductivity of the composite material, accelerated the charge transfer kinetics of the electrode, and improved the diffusion process of lithium ions by stabilizing the porous silicon microsphere structure. The lithium‑ion diffusion coefficient (D_Li⁺, cm²·s^-1) of the material can be calculated using the following equation^[49]:

  $ D_{\mathrm{Li}^{+}}=0.5\left[R T /\left(A n^2 F^2 \sigma_\text{w} c\right)\right]^2 $

  (3)

  Figure 8

  

  Figure 8.  (a) Nyquist plots of pSi and pSi/Sb-Sn@C electrodes after 300 cycles at 1 A·g^-1 (Inset: The corresponding equivalent circuit); (b) Linear fitting of Z′ vs ω^-0.5 plots in the low-frequency region

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  where R,T, A, n, F, σ_w, and c are the gas state constant (8.314 J·mol^-1·K^-1), thermodynamic temperature (K), electrode surface area (cm²), number of electrons transferred per mole of active substance (mol^-1), Faraday constant (9.648 5×10⁴ C·mol^-1), Warburg impedance coefficient (Ω·s^-0.5), and molar concentration of lithium ions (mol·L^-1), respectively. σ_w in Eq.3 can be derived from the following equation^[49]:

  $ Z^{\prime}=R_{\mathrm{s}}+R_{\mathrm{ct}}+\sigma_{\mathrm{w}} \omega^{-0.5} $

  (4)

  where Z′ is the real impedance part value of the Nyquist plot, and ω is the angular frequency (rad·s^-1). The linear relationship between the impedance real part values (Z′, Ω) of the pSi/Sb-Sn@C and pSi electrodes and ω^-0.5 in the low-frequency region is shown in Fig. 8b. σ_w may be determined by the slope of the straight line (Fig. 8b), and D_Li⁺ can be calculated in terms of Eq.3. The relevant results are also listed in Table 1. As shown in Table 1, compared to the pSi electrode, the pSi/Sb-Sn@C electrode exhibited lower ohmic polarization, faster electrochemical lithium storage kinetics, and enhanced lithium ion diffusion, which are the main reasons for its superior electrochemical lithium storage rate performance.

  Table 1

  

  Table 1.  Impedance parameter values obtained by the fitting of the Nyquist plots in Fig. 8a

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  Sample	Rs / Ω	Rct / Ω	CPE / μF	ZW / Ω	σ / (Ω·s-0.5)	DLi+ / (cm2·s-1)
  pSi	4.374	281.2	143.14	138.6	172.02	7.12×10-13
  pSi/Sb-Sn@C	2.199	49.55	385.37	59.46	129.15	1.26×10-12

  Compared with the recently reported silicon‑based anode materials for LIBs^{[8,21-22,50-55]}, the pSi/Sb-Sn@C composite prepared in this work exhibited enhanced electrochemical lithium storage performance, which is mainly attributed to its unique composition and microstructure. Surface deposition of Sb-Sn using porous silicon microspheres as a precursor and subsequent carbon coating result in the formation of pSi/Sb‑Sn@C composite with a stable 3D porous microsphere structure. The surface deposition of Sb/Sn and carbon coating significantly improved the electronic conductivity of the composite material, reduced the ohmic polarization of the electrode, and enhanced the mechanical stability of the 3D porous silicon structure. A stable 3D porous microsphere structure can increase the active sites for electrochemical lithium storage, reduce the diffusion distance of Li⁺, and buffer the significant volume changes of Si during the electrochemical lithium insertion/removal processes. The carbon coating can not only improve the electronic conductivity and mechanical stability of the composite but also avoid direct contact between silicon and electrolyte, thereby stabilizing the SEI film. In addition, Sb, Sn, and C as active lithium storage elements can contribute to a certain amount of electrochemical capacity. The above factors are responsible for the large reversible lithium storage capacity, excellent electrochemical lithium storage rate performance, and enhanced electrochemical charge‑discharge cycling stability of the pSi/Sb-Sn@C composite.

  3.   Conclusions

  Bimetallic (Sb-Sn) modified and carbon‑coated porous silicon microsphere composite (pSi/Sb‑Sn@C) has been successfully prepared by an experimental approach including alkaline etching and chemical displacement, heat treatment, acid etching, and carbon coating, using inexpensive industrial grade silicon aluminum alloy as raw material. The final product (pSi/Sb-Sn@C) had a core-shell structure as a whole. The 3D porous core pSi/Sb-Sn can increase the electrochemically active area, facilitate the diffusion of lithium ions, and buffer the huge volume changes of silicon during the charge-discharge processes. The carbon shell can avoid direct contact between the active components and the electrolyte, thereby stabilizing the SEI film. In addition, the introduction of active lithium storage components such as Sb, Sn, and C can improve the electronic conductivity of the silicon matrix, enhance the mechanical stability of porous silicon microspheres, and contribute to a certain amount of lithium storage capacity. Owing to its unique composition and microstructure, the constructed pSi/Sb-Sn@C composite exhibited excellent electrochemical lithium storage performance. This work utilizes the synergistic effect of multiple factors such as porous structure, active metal components, and carbon coating to improve the electrochemical lithium storage performance of silicon-based composite materials, which can provide a new strategy for the construction of high‑performance electrode materials for LIBs.

  
  Acknowledgments: This work was supported by the National Natural Science Foundation of China (Grant No.22372147).
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  Figure 1  Schematic illustration for the fabrication process of pSi and pSi/Sb-Sn@C composite

  In the naming of different materials, the symbol"/"represents the interface between the substrate and the metal deposition layer, and the symbol"-"is used to connect different components in the alloy or deposition layer.

  下载: 全尺寸图片 幻灯片
  

  Figure 2  (a) TEM image, (b) HRTEM images, and (c-h) EDS elemental mappings of pSi-Al/Zn-Sb-Sn

  下载: 全尺寸图片 幻灯片
  

  Figure 3  SEM images of SiAl alloy (a), pSi/Sb-Sn (b-d), and pSi/Sb-Sn@C (e, f)

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  Figure 4  (a) TEM image, (b) HRTEM image, and (c-h) EDS elemental mappings of pSi/Sb-Sn@align="center"

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  Figure 5  XRD patterns (a), enlarged XRD pattern (b), and Raman spectra (c) of pSi and pSi/Sb-Sn@C; TGA curves of pSi/Sb-Sn and pSi/Sb-Sn@C (d); N₂ adsorption-desorption isotherms of pSi (e) and pSi/Sb-Sn@C (f)

  Inset: the corresponding pore size distribution curves.

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  Figure 6  XPS spectra of the pSi/Sb-Sn@C composite: (a) full survey, (b) Si2p, (c) C1s, and (d) Sn3d; (e) Sb3d XPS spectra of pSi/Sb and pSi/Sb-Sn@C composites

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  Figure 7  (a) Initial CV curves at a scan rate of 0.1 mV·s^-1 and (b) galvanostatic charge-discharge curves at 100 mA·g^-1 of pSi/Sb-Sn@C electrode; (c, d) Rate performance of pSi and pSi/Sb-Sn@C electrodes; (e) Cycle performance curves of pSi, pSi/Sb-Sn, and pSi/Sb-Sn@C electrodes

  The bar chart represents the discharge capacity, while the line diagram represents the retention rate vs 3rd cycle in c.

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  Figure 8  (a) Nyquist plots of pSi and pSi/Sb-Sn@C electrodes after 300 cycles at 1 A·g^-1 (Inset: The corresponding equivalent circuit); (b) Linear fitting of Z′ vs ω^-0.5 plots in the low-frequency region

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  Table 1.  Impedance parameter values obtained by the fitting of the Nyquist plots in Fig. 8a

  Sample	Rs / Ω	Rct / Ω	CPE / μF	ZW / Ω	σ / (Ω·s-0.5)	DLi+ / (cm2·s-1)
  pSi	4.374	281.2	143.14	138.6	172.02	7.12×10-13
  pSi/Sb-Sn@C	2.199	49.55	385.37	59.46	129.15	1.26×10-12

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