English

• 0.   Introduction

  In the wake of the swift advancement of the new energy sector, lithium‑ion batteries (LIBs) have emerged as a pivotal energy alternative for portable electronic apparatuses, electric vehicles, and energy storage setups. This ascendancy can be chiefly attributed to their remarkable traits, namely, high energy density endowing them with prolonged power supply capabilities, an extended cycle life ensuring durability and reliability over numerous charge/discharge cycles, and a high operating voltage facilitating efficient energy conversion and utilization^[1-2]. However, the theoretical specific capacity of traditional graphite anode materials is only 372 mAh·g^-1, which is insufficient to meet the increasing demand for higher energy density^[3]. Therefore, exploring and developing high‑capacity, long-life anode materials has become one of the main focuses in LIB research^[4].

  Silicon (Si), endowed with a theoretical specific capacity of approximately 4 200 mAh·g^-1, has garnered significant attention from both academic and industrial communities as a prospective anode material, owing to its capacity to form lithium-silicon alloys (Li₂₂Si₅)^[5-6]. Nonetheless, during the charge and discharge processes, Si undergoes a volume expansion exceeding 300%, which triggers material pulverization, structural impairment, and continuous fracturing of the electrode‑ electrolyte interface. These adverse consequences severely undermine the cycling stability and lifetime of the battery^[7-8]. Consequently, notwithstanding their high theoretical capacity, Si anodes still encounter formidable technical hurdles in practical applications.

  In recent years, to address the volume expansion predicament of Si anodes, researchers have put forward the design concept of porous silicon and carbon composite materials^[9]. The fabrication of a porous structure can efficaciously cushion the volume expansion of Si during lithiation and delithiation procedures while concurrently furnishing supplementary channels that facilitate rapid lithium-ion transportation^[10]. Additionally, the integration of carbon materials not only augments the electrical conductivity of Si but also constructs a mechanically pliable framework, thereby further enhancing the cycling stability and electrochemical performance of the composite^[11]. Currently, the research efforts on porous silicon and carbon anodes predominantly center around structural design and optimization^[12-14], carbon coating modification and functionalization^[15-18], as well as the development and refinement of material preparation processes^[19-21]. However, merely depending on the active material structure design and single-layer carbon coating falls short of fulfilling the practical application requirements of porous silicon and carbon anode materials since issues such as material pulverization and deteriorated electrical contact induced by electrode expansion remain unresolved^[22]. In a recent approach, Li and his colleagues devised an innovative composite material, namely, double carbon shells coated with boron-doped porous Si (B-PSi@ double C). The synthesis process was ingeniously designed, commencing with the utilization of cost‑ effective industrial ferrosilicon alloys as the starting materials. Subsequently, a modified ball milling approach was implemented, which was then followed by the self-polymerization of dopamine hydrochloride. Finally, a chemical vapor deposition procedure with acetylene was carried out to complete the fabrication of this novel composite, opening up new possibilities for related research fields^[23]. Lim et al. addressed these issues by designing a porous Si structure combined with a double carbon-species coating layer, induced by low interfacial energy in a scalable process. Carbon and graphene were located on Si surfaces, forming a close interface that maintains electrical contact, suppressed lithium consumption, and enhanced charge transfer properties^[24].

  In light of this, we employed porous Si (pSi) derived from the dealloying reaction of silicon-aluminum alloys (Al₆₀Si₄₀) as the Si source, along with metal germanium (Ge) introduction and a dual-coating strategy of graphene (Gr) and carbon nanotubes (CNTs), to construct a pSi-based composite anode material (pSi/Ge@Gr/CNTs) with superior electrochemical performance. The introduction of metallic Ge atoms effectively modulated the mechanical properties of pSi, alleviated volume expansion, and enhanced the overall conductivity of the electrode material, thereby improving the battery′s overall performance. Half-cell tests with pSi/Ge@Gr/CNTs as the working electrode and lithium metal as the counter electrode showed that the specific capacity of the battery increased by more than 30% compared to a battery without Ge. Additionally, the cycling stability and rate performance of the battery were significantly improved. Furthermore, the dual wrapping of Gr and CNTs on the outer layer effectively enhanced the stability and conductivity of the electrode. Full-cell tests with pSi/Ge@Gr/CNTs as the anode and commercial LiNi_0.8Co_0.1Mn_0.1O₂ (NCM811) as the cathode further verified the practical potential of this electrode material. This work provides a novel approach to the design of high-performance hybrid pSi-based anode materials.

  1.   Experimental

  1.1   Preparation of pSi

  A total of 5 g of commercial Al₆₀Si₄₀ alloy powder, sourced from Changsha Tianjiu Metal Materials Co. Ltd., was gradually incorporated into 200 mL of hydrochloric acid solution with a concentration of 2 mol·L^-1. During this process, continuous stirring was maintained for 24 h to ensure thorough mixing and reaction. Afterward, the formed mixture was centrifuged and washed with deionized water and ethanol until neutral, then the bottom solid was dried in a vacuum oven at 60 ℃ for 8 h to obtain pSi. Meanwhile, we purchased nanosilicon (nSi, Shenzhen Kejing) with a particle size of 100 nm and conducted a comparative study with pSi.

  1.2   Synthesis of pSi/Ge

  The obtained pSi and Ge powder (Aladdin, 200 mesh, 99.999%) were mixed with a molar ratio of 9:1 and placed in a planetary corundum vacuum jar. The mass ratio of zirconia balls to the mixture was fixed at 12. The ball milling process involved 100 min grinding followed by a 10 min rest and repeated for a total of 10 cycles. The obtained sample was designated with pSi/Ge. Similarly, to determine the optimal ratio of pSi to Ge, we also prepared pSi/Ge with molar ratios of 12:1 and 1:1 for ball-milling, respectively. The samples were labeled as pSi₁/Ge₁ and pSi₁₂/Ge₁, respectively.

  1.3   Synthesis of pSi/Ge@Gr/CNTs

  First, graphene oxide (GO, Suzhou Carbon Tech) was washed by centrifugation until it became neutral. GO dispersion (mass fraction of 1%) was obtained after two hours of ultrasonic exfoliation treatment. 300 mg of pSi/Ge was dispersed in 40 mL of deionized water and sonicated for 20 min. Then, 1 mL of mass fraction of 35% poly(diallyldimethylammonium chloride) (PDDA, Aladdin) was added, and the mixture was subjected to sonication for another 10 min. After filtration to remove excess PDDA, the resulting mixture of pSi/Ge and PDDA (named pSi/Ge@PDDA) was uniformly dispersed into 100 mL of deionized water. To this solution, 40 mL of GO dispersion and 20 mL of carboxylated CNTs (XFNano) dispersion (mass fraction of 1%) were added, and the mixture was magnetically stirred to form a slurry with a certain viscosity. The slurry was heated at 80 ℃ until it solidified. The resulting solid powder of pSi/Ge@PDDA@GO/CNTs was ground with a mortar and placed in a horizontal tube furnace. Under a high-purity argon atmosphere, the sample was heated to 800 ℃ at a rate of 5 ℃·min^-1 and held for 2 h. After cooling to room temperature, the final product, pSi/Ge@Gr/CNTs, was obtained. As a comparison, we also prepared nSi/Ge@Gr/CNTs, pSi₁/Ge₁@Gr/CNTs, pSi₁₂/Ge₁@Gr/CNTs, pSi@Gr/CNTs, and pSi@CNTs without CNTs and Gr, respectively, following the same procedure.

  1.4   Material characterization

  A series of advanced analytical techniques were employed to characterize the samples. Powder X-ray diffraction (XRD) was carried out using a PANalytical instrument sourced from the Netherlands to investigate the crystal structure. XRD measurements were carried out at room temperature using Cu Kα radiation with a wavelength (λ) of 0.154 nm (with the voltage and current of 40 kV and 40 mA, as well as the scanning range of 5°-90°). In terms of visualizing the surface morphology, a field emission scanning electron microscope (SEM), specifically the Nova NanoSEM 450 model, was utilized. To further probe the morphology and microstructure at a finer scale, transmission electron microscopy (TEM), high-resolution transmission electron microscopy (HRTEM), and energy-dispersive X-ray spectroscopy (EDS) were conducted on a JEM-2100F apparatus, operating at an accelerating voltage of 200 kV. The surface chemical composition was analyzed via X-ray photoelectron spectroscopy (XPS) using the ESCALAB 250Xi system from Thermo Fisher Scientific. Before analyses, the samples were vacuum degassed at 150 ℃ for 10 h. Subsequently, N₂ adsorption‑desorption isotherms were acquired at 77 K with the help of a Micromeritics ASAP 2020 analyzer. To quantify the carbon content, thermogravimetric analysis (TGA) was performed using a STA449C thermal analyzer. The heating rate was set at 10 ℃·min^-1, ramping from room temperature to 800 ℃. Additionally, Raman spectra were collected using a Raman microscope (HORIBA Scientific LabRAM HR Evolution), with a 633 nm He-Ne laser serving as the light source for excitation. This multi-faceted analytical approach ensured a thorough understanding of the samples′ properties.

  1.5   Electrochemical testing

  The active component, conductive carbon black (Super-P), and polyvinylidene fluoride (PVDF) binder were mixed in N-methylpyrrolidone (NMP) with a mass ratio of 8:1:1 to form a slurry. This slurry was then meticulously applied onto copper foil and subjected to vacuum drying at 80 ℃ for 10 h. Once dried, the copper foil was precisely cut into circular electrodes with a diameter of 12 mm to be reserved for subsequent applications. The loading amount of composite materials in the electrodes was approximately 2.0 mg·cm^-2. For cell assembly, coin-type half-cells were constructed within an argon-purged glove box, where the oxygen and water vapor concentrations were strictly controlled below 0.1 mg·L^-1. A pure lithium foil was employed as the counter electrode, and a porous polypropylene membrane (Celgard 2400) served as the separator. The electrolyte consisted of 1 mol·L^-1 LiPF₆ dissolved in a 1:1 volume ratio of ethylene carbonate/diethyl carbonate (EC/DEC), supplemented with a mass fraction of 2% of vinylene carbonate (VC) as an additive.

  To evaluate the electrochemical performance, a series of tests were carried out. Constant current charge/discharge assays were executed using a Neware battery testing system, sourced from Shenzhen, China, within a voltage window of 0.01-2.0 V. Cyclic voltammetry (CV) measurements were conducted on a CHI660E electrochemical workstation, with the voltage sweep ranging from 0 to 3.0 V, while electrochemical impedance spectroscopy (EIS) was performed over a frequency range from 100 kHz to 0.01 Hz. Additionally, galvanostatic intermittent titration technique (GITT) measurements were obtained within the 0.01-2.0 V voltage range using a Neware system. Notably, the specific capacity of the half-cells was calculated concerning the total weight of the pSi/Ge@Gr/CNTs composite. The theoretical capacity of the pSi/Ge@Gr/CNTs material was calculated to be 2 238.04 mAh·g^-1 according to the proportion of each component. Additionally, the electrodes were pre-lithiated by directly contacting the pSi/Ge@Gr/CNTs electrode with lithium metal foil in the glove box for 15 min (with a certain amount of electrolyte between the two).

  1.6   Full cell assembly and performance evaluation

  The primary active component of the cathode was commercial NCM811, and the pre-lithiated pSi/Ge@Gr/CNTs electrode was used as the anode, with the negative and positive capacity ratio of approximately 1.1. The cathode was prepared by mixing commercial NCM811 powder, Super-P, and PVDF in a mass ratio of 8:1:1 in NMP to form a slurry. The same coating, drying, and slicing process was used for the cathode. Constant current charge/discharge experiments were carried out over a voltage interval spanning from 3.0 to 4.3 V. It should be emphasized that the specific capacity of these full cells was determined by considering the weight of the cathode material.

  2.   Results and discussion

  The preparation process of the pSi/Ge@Gr/CNTs composite material is illustrated in Fig. 1. First, commercial Al₆₀Si₄₀ alloy powder with a diameter of 1-3 μm (Fig.S1, Supporting information) was etched with hydrochloric acid at room temperature to remove Al from the Al₆₀Si₄₀ alloy. Subsequently, the resultant product underwent centrifugation. It was then thoroughly washed with deionized water and ethanol successively until the pH reached neutrality. Afterward, vacuum drying was performed to yield Si (Fig. 2a). Subsequently, the pSi and Ge were subjected to high-energy ball milling to introduce metallic Ge into pSi (pSi/Ge) (Fig. 2b). The pSi/Ge was then modified with PDDA to form a positive surface charge. The resulting pSi/Ge@PDDA was assembled with a two-dimensional Gr and one-dimensional CNTs network through electrostatic interactions to form pSi/Ge@PDDA@GO/CNTs. Finally, high-temperature annealing was conducted under an Ar atmosphere to prepare the pSi/Ge@Gr/CNTs composite material. For comparison, control samples were also prepared using the same process.

  Figure 1

  

  Figure 1.  Schematic illustration of the preparation of the pSi/Ge@Gr/CNTs composite material

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  Figure 2

  

  Figure 2.  SEM images of (a) pSi, (b) pSi/Ge, and (c) pSi/Ge@Gr/CNTs; (d) TEM and (e) HRTEM images of pSi/Ge@Gr/CNTs; (f-j) TEM image and elemental distribution maps of pSi/Ge@Gr/CNTs

  Inset: The lattice fringe images of the target selected region.

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  Fig.S2 shows SEM images and EDS spectra of the pSi/Ge sample. As seen from Fig. 2b and Fig.S2, the particle sizes of the pSi and pSi/Ge samples were similar, ranging from nanometers to micrometers. After ball milling, Si and Ge in the samples were uniformly mixed at the atomic level (Fig.S2c and S2d). The introduction of Ge improves the conductivity of the Si material, and the lattice mismatch between Si and Ge enhances ion migration rates. Fig. 2c and 2d display the SEM and TEM images of the pSi/Ge@Gr/CNTs sample, where Gr and CNTs were uniformly coated on the surface of the pSi/Ge. The interconnected Gr and CNTs form a three-dimensional hybrid network structure, which buffers the volume changes of Si during the charge/discharge process and provides pathways for rapid charge transport. In the HRTEM images, the measured lattice spacings of 0.31 and 0.33 nm correspond to the Si(111) and Ge(111) crystal planes^[25-26], respectively (Fig. 2e, S3b, and S3c). Additionally, the element distribution maps (Fig. 2f-2j) and EDS spectra (Fig.S3) showed that Ge was uniformly distributed within the Si and encapsulated by carbon materials, with no other impurity elements detected. These characterizations confirm that the pSi/Ge@Gr/CNTs sample consisted of Ge-introduced pSi wrapped with a hybrid layer of Gr and CNTs.

  The phase structure of the prepared samples was characterized by XRD, as shown in Fig. 3a and S4. The XRD patterns of the Al₆₀Si₄₀ alloy and pSi are presented in Fig.S4a. The diffraction peaks of the obtained pSi matched the standard card for elemental Si (PDF No.27-1402), and no diffraction peaks for Al were detected, indicating that the aluminum in the alloy was completely removed by hydrochloric acid etching. After ball milling with Ge, both Si and Ge diffraction peaks were detected in the pSi/Ge sample, as shown in Fig. 3a. In the pSi/Ge@Gr/CNTs sample, the intensity of the Si diffraction peaks was significantly weaker than that of pSi/Ge, and the characteristic peaks of Ge were nearly absent, suggesting that the signals of Si and Ge in pSi/Ge@Gr/CNTs were shielded by the three-dimensional network structure composed of Gr and CNTs, further confirming the integrity of the hybrid layer coating. Additionally, the pSi@CNTs sample showed the characteristic diffraction peaks of CNTs, as shown in Fig.S4b, due to the introduction of only CNTs. Fig. 3b shows the Raman spectra of the pSi/Ge and pSi/Ge@Gr/CNTs samples. The two peaks at 511 and 948 cm^-1 in both samples correspond to crystalline Si, while the Raman scattering at 300 cm^-1 is attributed to metallic Ge^[27-28]. Furthermore, the two peaks at 1 352 and 1 584 cm^-1 in the pSi/Ge@Gr/CNTs sample correspond to the D and G bands of carbon materials, respectively, which are caused by the coating of Gr and CNTs^[29-30].

  Figure 3

  

  Figure 3.  (a) XRD patterns of pSi/Ge and pSi/Ge@Gr/CNTs; (b) Raman spectra of pSi, pSi/Ge, and pSi/Ge@Gr/CNTs; (c) C1s, (d) Si2p, (e) O1s, and (f) Ge3d XPS spectra of pSi/Ge@Gr/CNTs

  Inset: Raman spectra of the local amplification of pSi and pSi/Ge.

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  Furthermore, it should be noted that, as shown in the insert in Fig. 3b, after Ge was introduced into pSi, the Raman characteristic peak of Si—Si at approximately 511 cm^-1 exhibited a shift towards lower wavenumbers (around 500 cm^-1). The main reason might be that the atomic radius of Ge is larger than that of Si. When Ge atoms are introduced into the Si lattice, the Si lattice will expand and distort, resulting in changes in the length of Si—Si bonds. According to the relationship between the vibration frequency of Raman spectra and the bond force constant as well as the reduced mass, changes in bond length will lead to variations in the bond force constant, thereby causing shifts in the positions of Raman peaks.

  The elemental composition and valence states of the sample surfaces were characterized using XPS. Fig.S5 shows the XPS survey and high-resolution spectra of the pSi/Ge sample, with its main components being Si, Ge, and O elements. The oxygen mainly originates from the oxidation of Si and Ge, which is consistent with the Raman results. The XPS survey spectrum in Fig.S6 confirms the presence of Si, Ge, C, and O elements in the pSi/Ge@Gr/CNTs sample. Fig. 3c shows the high-resolution C1s XPS spectrum of the pSi/Ge@Gr/CNTs sample, where the deconvoluted peaks correspond to C—C, C—O—C, and C=O bonds at 284.8, 286.1, and 289.3 eV, respectively^[31]. Fig. 3d shows the high-resolution Si2p XPS spectrum, with peaks at 99.5, 100.8, and 103.8 eV, corresponding to Si—Si, Si—O—C, and Si—O bonds, respectively^[32]. The presence of the Si—O—C bond confirms the tight interconnection between Si and the conductive network of Gr and CNTs, which promotes electron transfer and buffers the volume changes of Si. In the high-resolution O1s XPS spectrum (Fig. 3e), the deconvoluted peak at 533.1 eV is attributed to Si—O and C—O bonds, while the peak at 531.5 eV corresponds to C=O bonds^[33-35]. Fig. 3f shows the high-resolution Ge3d spectrum, with peaks at 28.4 and 33.3 eV corresponding to Ge—Ge and Ge—O bonds, respectively^[36].

  Fig. 4a and 4b display the N₂ adsorption-desorption isotherm and pore size distribution of pSi/Ge@Gr/CNTs. The sample exhibited typical type Ⅳ isotherms and H3 hysteresis loops, confirming the presence of mesoporous structures in the composite, with a specific surface area of 74 m²·g^-1 and an average pore size of 20.23 nm. The abundant porous structure facilitates contact between the electrolyte and active sites while accommodating the volume changes of Si during charge/discharge processes. To determine the content of Gr and CNTs in the samples, TGA was performed. As shown in Fig. 4c and S7, the mass fraction of carbon in the pSi/Ge@Gr/CNTs, pSi@Gr/CNTs, and pSi@CNTs samples were approximately 54%, 48%, and 58%, respectively.

  Figure 4

  

  Figure 4.  (a) N₂ adsorption-desorption isotherm, (b) pore distribution, and (c) TGA curve of pSi/Ge@Gr/CNTs

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  The lithium storage performance of the prepared pSi/Ge@Gr/CNTs electrode was further evaluated through electrochemical tests. Fig. 5a shows the CV curves of the pSi/Ge@Gr/CNTs electrode for the first three cycles within a potential window of 0-3 V (scanning rate of 0.2 mV·s^-1). During the first discharge, the broad peak around 0.80-1.70 V represents the formation of the solid electrolyte interface (SEI) layer, while the sharp peak below 0.10 V corresponds to the reaction of crystalline Si with lithium to form amorphous Li_xSi alloy^[37-38]. During the subsequent charge, a reduction peak at 0.20 V is related to the amorphization of Si, and two oxidation peaks at 0.34 and 0.50 V correspond to the de-alloying reaction of the Li-Si alloy^[39]. Compared to the CV curves of pSi@Gr/CNTs in Fig.S8, the pSi/Ge@Gr/CNTs electrode exhibited less polarization during the second and third cycles, with better reversibility and sharper peaks, indicating faster reaction kinetics.

  Figure 5

  

  Figure 5.  (a) CV curves of the pSi/Ge@Gr/CNTs electrode during the first three cycles at a scan rate of 0.2 mV·s^-1; (b) Initial charge/discharge curves of the pSi/Ge@Gr/CNTs electrode before and after pre-lithiation treatment; (c) Charge/discharge curves of the pSi/Ge@Gr/CNTs electrode at different cycle numbers; (d) Rate performance of pSi@CNTs, pSi@Gr/CNTs, and pSi/Ge@Gr/CNTs electrodes; Cycling performance of pSi@CNTs, pSi@Gr/CNTs, and pSi/Ge@Gr/CNTs at (e) 0.2 A·g^-1 and (f) 1 A·g^-1

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  Fig. 5b shows the first constant current charge/ discharge curves of the pSi/Ge@Gr/CNTs electrode before and after pre-lithiation, at a current density of 0.2 A·g^-1 and a voltage range of 0.01 to 2 V (vs Li/Li⁺). Without pre-lithiation, the initial charge and discharge capacities of the pSi/Ge@Gr/CNTs electrode were 1 162 and 2 005 mAh·g^-1, respectively, corresponding to an initial Coulombic efficiency (ICE) of 58%. The significant irreversible capacity loss may be due to the formation of an irreversible SEI film on the electrode surface during the first discharge. After pre-lithiation, the initial charge and discharge capacities were 1 194 and 1 258 mAh·g^-1, respectively, with an ICE of 95%. Fig. 5c displays the constant current charge/discharge curves of the pSi/Ge@Gr/CNTs electrode during the 1st, 10th, 20th, 50th, and 100th cycles, showing good cycling stability, with a capacity retention of 66% after 100 cycles.

  Fig. 5d presents the rate capability of different electrodes. At current densities of 0.2, 0.5, 0.8, 1.0, 2.0, and 3.0 A·g^-1, the average specific discharge capacities of the pSi/Ge@Gr/CNTs electrode were 700, 562, 506, 467, 379, and 286 mAh·g^-1, respectively. Furthermore, when the current density returned to 0.2 A·g^-1, the reversible specific capacity recovered to 650 mAh·g^-1. In contrast, the pSi@Gr/CNTs and pSi@CNTs electrodes exhibited lower capacities at the same current densities, indicating that the pSi/Ge@Gr/CNTs electrode had excellent high‑rate charge/discharge capabilities. It should be pointed out that the battery for testing the rate performance in Fig. 5d had been pre-cycled for 100 cycles at a current density of 0.2 mA·g^-1. Therefore, the reversible capacity at a current density of 0.2 mA·g^-1 in Fig. 5d was lower than the initial capacity in Fig. 5e. Fig. 5e, S9, and 5f show the cycling stability of pSi/Ge@Gr/CNTs at current densities of 0.2, 0.5, and 1.0 A·g^-1. Compared to the pSi@Gr/CNTs and pSi@CNTs electrodes, the pSi/Ge@Gr/CNTs electrode demonstrated superior specific capacity and cycling stability. This is attributed to the synergistic effect of Ge introduction and the hybrid coating layer, which improves the mechanical and electrochemical performance of pSi. It can be seen from some comparative data listed in Table S1 that the pSi/Ge@Gr/CNTs electrode is competitive in electrochemical performance. Figs.S9b and S9c present the cycling performance and rate performance curves of pSi₁/Ge₁@Gr/CNTs, pSi₁₂/Ge₁@Gr/CNTs, and nSi/Ge@Gr/CNTs electrodes at a current density of 1.0 A·g^-1. From the figure, it can be observed that as the Ge content increased, the specific capacity of the electrodes showed an upward trend. However, considering that the theoretical capacity of Ge was significantly lower than that of Si, using the pSi₁/Ge₁@Gr/CNTs material with a pSi-Ge molar ratio of 1:1 as the electrode did not result in the optimal specific capacity. When the pSi/Ge@Gr/CNTs material with a pSi-Ge molar ratio of 9:1 was used as the electrode, it exhibited the most excellent electrochemical performance, indicating that a pSi-Ge molar ratio of 9:1 was the optimal ratio. In addition, we also carried out similar composite treatments using nSi with a particle size of 100 nm and conducted electrochemical tests. The results clearly showed that the materials containing pSi were significantly superior to nSi in terms of both rate performance and long-cycle stability, further confirming that pSi can effectively alleviate the volume expansion problem and thus improve the electrochemical performance of the materials.

  Fig. 6a and S10 show the EIS of the pSi@Gr/CNTs and pSi/Ge@Gr/CNTs electrodes after different numbers of cycles (the corresponding equivalent circuit diagrams are shown in the insets, where R_s stands for solution resistance, R_ct stands for charge transfer resistance, R_SEI stands for resistance of SEI film, Z_W stands for diffusion resistance, CPE₁ stands for electric double layer capacitance, CPE₂ stands for capacitance of SEI film). There were significant differences in the EIS before and after cycling due to the formation of the SEI film and the increased electron transfer rate. Fig. 6a shows the EIS of the two electrodes after five cycles, where the semicircle in the high-frequency region corresponds to the R_SEI, and the semicircle in the mid-frequency region corresponds to the R_ct^[40-41]. Both the R_SEI and R_ct of the pSi/Ge@Gr/CNTs electrode were smaller than those of the pSi@Gr/CNTs electrode, indicating that Ge introduction reduces the electrode resistance and improves its conductivity. After 50 cycles, the EIS of the pSi/Ge@Gr/CNTs electrode remained relatively stable, proving that the electrode maintains structural stability during cycling (Fig.S10).

  Figure 6

  

  Figure 6.  (a) EIS of the pSi@Gr/CNTs and pSi/Ge@Gr/CNTs electrodes after 5 cycles; (b) Schematic illustration of the structural evolution of pSi and pSi/Ge@Gr/CNTs during cycling

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  Fig. 6b illustrates the structural evolution of pSi and pSi/Ge@Gr/CNTs during cycling. During cycling, pSi particles underwent fracture, and the freshly exposed Si surface continuously reacted with the electrolyte, forming a thick SEI film, leading to battery failure. In contrast, the pSi/Ge@Gr/CNTs electrode, enhanced by introducing Ge, exhibited improved fracture toughness. The Gr and CNT network coating, along with the porous structure, mitigated the expansion and fracture of Si particles, forming an effective SEI film on the surface that minimizes irreversible lithium-ion loss, resulting in rapid charge transfer kinetics. Additionally, as seen in the SEM and sectional SEM images of the pSi/Ge@Gr/CNTs electrode before and after 50 cycles (Fig.S11), no significant cracks were observed, the volume expansion rate was low, and the structure remained intact, further supporting the above conclusions.

  To investigate the reaction kinetics and lithium storage mechanism of the pSi/Ge@Gr/CNTs electrode, CV tests were conducted at different scan rates (Fig. 7a). The lithium-ion storage in electrode materials can be divided into diffusion-controlled charge storage and pseudocapacitance-controlled charge storage. The respective capacity contributions were calculated using the following equation^[42]:

  $ i_{\mathrm{p}}=a v^{\mathrm{b}} $

  (1)

  Figure 7

  

  Figure 7.  (a) CV curves of the pSi/Ge@Gr/CNTs electrode at different scan rates; (b) Linear relationship between lg i_p and lg v for the oxidation and reduction peaks of the pSi/Ge@Gr/CNTs electrode; (c) Contribution ratio of capacitance from pseudocapacitance and diffusion-controlled region of the pSi/Ge@Gr/CNTs electrode at 0.8 mV·s-1; (d) Contribution ratio at various scan rates of the pSi/Ge@Gr/CNTs electrode; (e, f) Li⁺ diffusion coefficients of pSi@Gr/CNTs and pSi/Ge@Gr/CNTs electrodes during the charge/discharge process

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  where v is the scan rate and i_p is the peak current. The value of b can be determined by fitting the lg i_p and lg v curves. When b=0.5, the charge storage mechanism is diffusion-controlled, and when b=1, it is pseudocapacitance-controlled^[43]. Fig. 7b shows the fitting curves of the oxidation and reduction peaks of the pSi/Ge@Gr/CNTs electrode, with slopes of 0.52 and 0.57, respectively. This indicates that the lithium storage mechanism in the pSi/Ge@Gr/CNTs electrode is primarily diffusion-controlled.

  Moreover, the contribution of pseudocapacitance to the diffusion-controlled process can be quantified using the following equations^[44]:

  $ i=k_1 v+k_2 v^{1 / 2} $

  (2)

  $ i / v^{1 / 2}=k_1 v^{1 / 2}+k_2 $

  (3)

  where k₁v and k₂v^1/2 represent the contributions from pseudocapacitance and diffusion‑controlled mechanisms, respectively. Fig. 7c and S12 show the separation of the pseudocapacitance-controlled and diffusion-controlled portions from the CV curves at different scan rates. As illustrated in Fig. 7d, with the increase in scan rate, the impact of capacitive control became more pronounced, and when the scan rate reached 1.2 mV·s^-1, the capacitive contribution reached as high as 72%.

  Additionally, we measured the lithium-ion diffusion capacity of the pSi@Gr/CNTs and pSi/Ge@Gr/CNTs electrodes using GITT tests (Fig.S14). According to Fick′s second law of diffusion, the lithium-ion diffusion coefficient (D) was calculated using the following simplified equation^[45]:

  $ D_{\mathrm{GITT}}=\frac{4}{\pi \tau} \times\left(\frac{m_{\mathrm{B}} V_{\mathrm{m}}}{M_{\mathrm{B}} S}\right)^2 \times\left(\frac{\Delta E_{\mathrm{s}}}{\Delta E_\tau}\right)^2 $

  (4)

  where τ is the duration of the current pulse, m_B and M_B are the mass and molar mass of the material, S is the electrode surface area, V_m is the molar volume, and ΔE_S and ΔE_τ are the voltage changes between steps and during the current pulse, respectively. Fig. 7e and 7f show the Li⁺ diffusion coefficients of pSi@Gr/CNTs and pSi/Ge@Gr/CNTs during charge/discharge. The diffusion coefficient of pSi/Ge@Gr/CNTs was higher than that of pSi@Gr/CNTs, indicating that the Ge introduction effectively enhances lithium-ion diffusion.

  To further assess the practical application prospects of the pSi/Ge@Gr/CNTs electrode, a coin full-cell was assembled with commercial NCM811 as the cathode and pSi/Ge@Gr/CNTs as the anode. As shown in Fig. 8a, the NCM811||pSi/Ge@Gr/CNTs full cell retained a specific capacity of 80 mAh·g^-1 after 100 cycles at a 0.2 A·g^-1 current density. The charge/ discharge curves under different cycles in Fig. 8b demonstrated the good operational stability of the NCM811||pSi/Ge@Gr/CNTs full cell. Lastly, as shown in Fig. 8c, the ability of the full cell to power an LED (light-emitting diode) array further proves its potential application performance.

  Figure 8

  

  Figure 8.  (a) Cycling performance of the NCM811||pSi/Ge@Gr/CNTs full cell at a current density of 0.2 A·g^-1; (b) Charge/discharge curves of the NCM811||pSi/Ge@Gr/CNTs full cell at different cycle numbers; (c) LED light powered by the full cell

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  3.   Conclusions

  In conclusion, we developed a Si-based composite anode material (pSi/Ge@Gr/CNTs) with excellent electrochemical performance by using porous silicon (pSi), derived from the dealloying of an Al₆₀Si₄₀ alloy, as the silicon source and introducing Ge into it. This material was further enhanced with a hybrid-coating strategy involving graphene (Gr) and carbon nanotubes (CNTs). The pSi/Ge@Gr/CNTs anode material for LIBs exhibited the following key advantages: (ⅰ) The voids between the porous structures provide extra space to accommodate the volume expansion of Si during charge/ discharge cycles. Additionally, the porous structure enhances lithium-ion transport and electrolyte infiltration, ensuring good contact between the electrolyte and the anode. (ⅱ) Ge introduction improves the fracture toughness of the Si-based material, significantly enhancing the electrical conductivity of the electrode. This accelerates the transport of lithium ions and electrons, resulting in higher specific capacity and extended cycling life. (ⅲ) The combination of Gr and CNTs forms a stable network structure, providing high conductivity and flexibility. This network also acts as a protective layer, buffering the volume expansion of Si. We believe that this work offers new insights into the design and modification of pSi anodes.

  
  Supporting information is available at http://www.wjhxxb.cn
  Conflicts of interest: There are no conflicts to declare.
  
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• 

  Figure 1  Schematic illustration of the preparation of the pSi/Ge@Gr/CNTs composite material

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  Figure 2  SEM images of (a) pSi, (b) pSi/Ge, and (c) pSi/Ge@Gr/CNTs; (d) TEM and (e) HRTEM images of pSi/Ge@Gr/CNTs; (f-j) TEM image and elemental distribution maps of pSi/Ge@Gr/CNTs

  Inset: The lattice fringe images of the target selected region.

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  Figure 3  (a) XRD patterns of pSi/Ge and pSi/Ge@Gr/CNTs; (b) Raman spectra of pSi, pSi/Ge, and pSi/Ge@Gr/CNTs; (c) C1s, (d) Si2p, (e) O1s, and (f) Ge3d XPS spectra of pSi/Ge@Gr/CNTs

  Inset: Raman spectra of the local amplification of pSi and pSi/Ge.

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  Figure 4  (a) N₂ adsorption-desorption isotherm, (b) pore distribution, and (c) TGA curve of pSi/Ge@Gr/CNTs

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  Figure 5  (a) CV curves of the pSi/Ge@Gr/CNTs electrode during the first three cycles at a scan rate of 0.2 mV·s^-1; (b) Initial charge/discharge curves of the pSi/Ge@Gr/CNTs electrode before and after pre-lithiation treatment; (c) Charge/discharge curves of the pSi/Ge@Gr/CNTs electrode at different cycle numbers; (d) Rate performance of pSi@CNTs, pSi@Gr/CNTs, and pSi/Ge@Gr/CNTs electrodes; Cycling performance of pSi@CNTs, pSi@Gr/CNTs, and pSi/Ge@Gr/CNTs at (e) 0.2 A·g^-1 and (f) 1 A·g^-1

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  Figure 6  (a) EIS of the pSi@Gr/CNTs and pSi/Ge@Gr/CNTs electrodes after 5 cycles; (b) Schematic illustration of the structural evolution of pSi and pSi/Ge@Gr/CNTs during cycling

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  Figure 7  (a) CV curves of the pSi/Ge@Gr/CNTs electrode at different scan rates; (b) Linear relationship between lg i_p and lg v for the oxidation and reduction peaks of the pSi/Ge@Gr/CNTs electrode; (c) Contribution ratio of capacitance from pseudocapacitance and diffusion-controlled region of the pSi/Ge@Gr/CNTs electrode at 0.8 mV·s-1; (d) Contribution ratio at various scan rates of the pSi/Ge@Gr/CNTs electrode; (e, f) Li⁺ diffusion coefficients of pSi@Gr/CNTs and pSi/Ge@Gr/CNTs electrodes during the charge/discharge process

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  Figure 8  (a) Cycling performance of the NCM811||pSi/Ge@Gr/CNTs full cell at a current density of 0.2 A·g^-1; (b) Charge/discharge curves of the NCM811||pSi/Ge@Gr/CNTs full cell at different cycle numbers; (c) LED light powered by the full cell

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