In this review, I systematically examine the recent progress in porous silicon–carbon composite anodes for lithium-ion cells. The rapid expansion of electric vehicles and grid-scale energy storage has placed unprecedented demands on the energy density, rate capability, and cycle life of lithium-ion cells. Conventional graphite anodes, although mature and reliable, are constrained by a theoretical specific capacity of only 372 mAh·g−1. Silicon-based materials have attracted intense attention because of their extremely high theoretical specific capacity, moderate lithium insertion potential, and abundant availability. However, the practical implementation of silicon in lithium-ion cells remains hindered by severe volume changes during alloying and dealloying, poor electronic conductivity, and continuous solid electrolyte interphase growth. Porous architectures combined with carbon matrices have therefore emerged as one of the most promising strategies to address these coupled challenges.
The theoretical capacity of silicon can be estimated from the fully lithiated phase according to the following relation:
$$ C_{th} = \frac{nF}{3.6M} $$
where \(n\) is the number of electrons transferred per silicon atom, \(F\) is the Faraday constant, and \(M\) is the molar mass. For the formation of \(\mathrm{Li_{4.4}Si}\), \(n \approx 4.4\), \(F = 96485 \ \mathrm{C \cdot mol^{-1}}\), and \(M = 28.09 \ \mathrm{g \cdot mol^{-1}}\), giving a value close to 4200 mAh·g−1. This value is approximately ten times that of graphite. Nevertheless, the volume expansion during lithiation can exceed 300%:
$$ \frac{\Delta V}{V_0} = \frac{V_{\mathrm{Li_xSi}} – V_{\mathrm{Si}}}{V_{\mathrm{Si}}} \times 100\% \approx 300\% $$
Such a large volumetric change generates substantial mechanical stress, which can be approximated by a simple elastic relation:
$$ \sigma \propto E \frac{\Delta V}{V_0} $$
where \(E\) is the Young’s modulus of the surrounding matrix. When this stress exceeds the fracture strength of the electrode components, particle pulverization, electrode delamination, and continuous solid electrolyte interphase reconstruction occur. The repeated rupture and regeneration of the solid electrolyte interphase consumes active lithium and electrolyte, leading to low initial Coulombic efficiency, rapid capacity fading, and poor rate performance in lithium-ion cells.
In my analysis, neither porosity alone nor carbon coating alone can fully resolve the coupled mechanical and electrochemical degradation of silicon anodes. The synergistic construction of porous silicon and carbon matrices is therefore essential. The core design principle is to coordinate multiple conflicting requirements: the porous structure must buffer volume strain, while the continuous carbon network must provide electronic conduction; internal voids must accommodate expansion, while external constraints must maintain structural integrity. Based on spatial configuration, I classify porous silicon–carbon composites into three major families: core–shell structures, yolk–shell structures, and embedded/encapsulated structures. The following sections discuss their preparation methods, electrochemical performance, failure mechanisms, and application suitability.
1. Preparation Methods for Porous Silicon–Carbon Composites
The performance of porous silicon–carbon composites is determined not only by their composition but also by the precise control of pore architecture and silicon–carbon interfaces. In this section, I examine the preparation of porous silicon skeletons and the subsequent strategies for silicon–carbon integration. The preparation routes are summarized in Table 1 and Table 2.
1.1 Preparation of Porous Silicon
The porous silicon skeleton is the central structural unit that governs carrier transport, strain tolerance, and interfacial stability. Therefore, the fabrication of porous silicon with suitable pore size, pore volume, and pore connectivity is the physical foundation for high-performance composite anodes in lithium-ion cells. The main preparation methods include template methods, magnesiothermic reduction, and dealloying/etching.
1.1.1 Template Method
The template method constructs porous silicon skeletons by using sacrificial templates. Its core advantage is the precise control of pore size, geometry, and spatial arrangement, which provides an ideal platform for studying structure–performance relationships. Depending on the template nature, hard-template and soft-template routes are distinguished.
Hard-template methods typically employ rigid templates such as silica microspheres or metal–organic frameworks. For example, recycled photovoltaic silicon waste can be combined with calcium carbonate as a hard template to produce porous silicon–carbon composites. Such materials have demonstrated a reversible specific capacity of approximately 946.3 mAh·g−1 after 200 cycles at 500 mA·g−1. In contrast, soft-template methods are generally milder and the templates are easier to remove. A typical soft-template route uses a triblock copolymer to synthesize mesoporous silica, which is then converted into porous silicon with endogenous carbon. The resulting composite can deliver an initial reversible capacity as high as 3058 mAh·g−1.
The template method offers exceptional control over porosity. The pore volume fraction can be expressed as:
$$ P = \left(1 – \frac{\rho_{app}}{\rho_{th}}\right) \times 100\% $$
where \(\rho_{app}\) is the apparent density and \(\rho_{th}\) is the theoretical density. The specific surface area can be estimated from nitrogen adsorption isotherms using the Brunauer–Emmett–Teller equation:
$$ S_{BET} = \frac{V_m N_A \sigma_0}{22414} $$
where \(V_m\) is the monolayer adsorbed volume, \(N_A\) is Avogadro’s number, and \(\sigma_0\) is the cross-sectional area of the adsorbate molecule. Despite these advantages, template methods usually involve template synthesis, precursor infiltration, and template removal. The process is complex, time-consuming, and often expensive. Moreover, some etching agents pose safety and environmental risks, which limits large-scale production for lithium-ion cells.
1.1.2 Magnesiothermic Reduction
Magnesiothermic reduction is a classic thermochemical route in which magnesium reduces silica to silicon and magnesium oxide. After acid washing to remove MgO, a porous silicon skeleton is obtained. The overall reaction can be written as:
$$ 2\mathrm{Mg} + \mathrm{SiO_2} \rightarrow 2\mathrm{MgO} + \mathrm{Si} $$
This method is attractive because of its scalability and relatively low cost. For instance, spherical porous silicon has been prepared from microsilica fume generated by the photovoltaic industry. The resulting composite retained a reversible capacity of about 803.92 mAh·g−1 after 300 cycles at 1000 mA·g−1. In another approach, commercial silica sol and a resorcinol–formaldehyde resin were spray-dried and calcined to form silica microspheres, followed by magnesiothermic reduction to produce \(\mathrm{M-pSi@C}\). This anode exhibited a capacity retention of 95% after 500 cycles at 2 A·g−1.
However, magnesiothermic reduction has limited ability to precisely control pore architecture. The pores mainly originate from the space left by MgO byproducts after etching, and their uniformity and ordering are inferior to those obtained by template methods. The strongly exothermic reaction can also cause excessive growth of silicon crystallites, reducing surface area and electrochemical activity. Nevertheless, because of its simple process, low cost, and easy scale-up, magnesiothermic reduction remains one of the most industrially promising routes for porous silicon in lithium-ion cells.
1.1.3 Dealloying and Etching
Dealloying and etching methods obtain porous silicon by selectively removing components from a parent alloy or compound. Their advantages include wide raw-material adaptability and high green-recycling value, but the pore structure strongly depends on the precursor composition, making independent design and precise control difficult.
In a typical dealloying process, an Al–Si alloy powder is prepared by gas atomization and then etched with hydrochloric acid to remove the aluminum matrix. The resulting three-dimensional interconnected porous silicon skeleton can maintain a reversible capacity of about 1607 mAh·g−1 after 200 cycles at 0.1 C. Beyond acid etching, other chemical pathways can also produce porous silicon. For example, \(\mathrm{Mg_2Si}\) and graphitic carbon nitride can be subjected to two-step heat treatment and vacuum volatilization to obtain porous silicon in one step. The composite retained 95.68% of its capacity after 4000 cycles at 5 A·g−1.
Dealloying/etching combines cost advantages with good scalability, showing significant potential for large-scale porous silicon production. However, pore uniformity and precise controllability are inferior to template methods. Future improvements may include precursor engineering, multistage or selective etching, and green etching with waste-liquid recycling. The detailed comparison of porous silicon preparation methods is presented in Table 1.
| Method | Pore control | Cost | Scalability | Main advantage | Main challenge |
|---|---|---|---|---|---|
| Template method | High | High | Low | Precise pore size and geometry | Complex template removal and high cost |
| Magnesiothermic reduction | Moderate | Moderate | High | Simple process and low raw-material cost | Byproduct removal and limited pore ordering |
| Dealloying/etching | Moderate to low | Low | High | Good recycling value and simple equipment | Strong acid waste and nonuniform pores |
1.2 Preparation of Silicon–Carbon Composites
After constructing the porous silicon skeleton, the next critical step is to integrate silicon with a carbon matrix. According to the carbon introduction route and the silicon–carbon interaction, I classify the composite strategies into chemical vapor deposition, liquid impregnation and pyrolysis, and ball milling. Each strategy influences the electronic conductivity, mechanical buffering, and interfacial stability of the final anode in lithium-ion cells.
1.2.1 Chemical Vapor Deposition
Chemical vapor deposition uses gaseous precursors to deposit solid films on substrates. Typically, hydrocarbons are used as carbon sources and are decomposed at high temperature, allowing active carbon atoms to deposit on porous silicon surfaces and form a continuous dense carbon layer. This method enables atomic-level interface engineering and is ideal for constructing core–shell structures.
For example, porous silicon microspheres composited with carbon nanotubes can be coated with a uniform carbon layer of about 4 nm using acetylene at 700 °C for 30 min. The resulting \(\mathrm{CNT@Si@C}\) composite retained 87% of its capacity after 1500 cycles. In another study, ethylene was used at 1080 °C to grow nitrogen-doped vertical graphene nanosheets on a silicon/porous carbon fiber substrate, producing a binder-free composite anode with a three-dimensional continuous conductive network. This material retained 83.5% of its capacity after 3000 cycles at 1.0 A·g−1.
The carbon deposition rate in CVD can be described by an Arrhenius-type relation:
$$ r_{dep} = A \exp\left(-\frac{E_a}{RT}\right) $$
where \(A\) is a pre-exponential factor, \(E_a\) is the activation energy, \(R\) is the gas constant, and \(T\) is the absolute temperature. Although CVD provides unique advantages in interface control, it suffers from complex equipment, high energy consumption, slow deposition, and high cost. For three-dimensional porous structures, uneven coating and discontinuous conductive networks can also occur, making scale-up difficult for practical lithium-ion cells.
1.2.2 Liquid Impregnation and Pyrolysis
Liquid impregnation and pyrolysis involve infiltrating carbon precursors such as pitch or phenolic resin into a porous silicon skeleton. Capillary action fills the pores, and subsequent evaporation, crosslinking, and high-temperature carbonization convert the precursor into a conductive carbon layer. This method builds a three-dimensional continuous conductive network inside and outside the silicon skeleton.
For instance, petroleum pitch can be infiltrated into nanoporous silicon under vacuum, followed by carbonization to form a continuous carbon matrix. When combined with graphite, the composite anode in a full lithium-ion cell with an NMC532 cathode retained more than 80% of its capacity after 450 cycles at 3.3 mAh·cm−2. To further balance structural stability and kinetics, a blend of asphaltene and resin was impregnated into porous silicon and carbonized at 700 °C for 2 h, producing a defect-rich continuous carbon network. This \(\mathrm{P-Si@DCN}\) composite retained 85% of its capacity after 300 cycles at 1 A·g−1.
Liquid impregnation and pyrolysis offer prominent advantages in performance, cost, and scalability, making them suitable for constructing three-dimensional composite structures from high-porosity silicon. However, carbon layer uniformity, mass control, and interfacial stability still require optimization. Future improvements may involve multicomponent formulations, interface modification, and combination with CVD to enhance overall performance in lithium-ion cells.
1.2.3 Ball Milling
Ball milling is a physical composite method that uses high-energy mechanical forces to achieve mixing, fracturing, cold welding, and mild chemical reactions among solid powders. When preparing silicon–carbon composites, porous silicon and carbon materials such as graphite or carbon black are placed in a ball-milling jar. The grinding balls create intimate contact and uniform dispersion at the microscale. The method is simple, efficient, and easy to scale up.
For example, nanosilicon and tin dioxide can be ball-milled to form a \(\mathrm{Si@SnO_2}\) intermediate, followed by secondary ball milling with nitrogen/phosphorus co-doped carbon to construct a sandwich-structured \(\mathrm{C/Si@SnO_2}\) anode. This material retained 90% of its capacity after 1000 cycles at 2 A·g−1. In another route, silicon cutting waste from the photovoltaic industry was mixed with graphite, pitch, and polyvinylpyrrolidone, then high-energy ball-milled to produce a nitrogen-doped \(\mathrm{Si/G@voids@C}\) composite with porous structure. The first discharge specific capacity reached 2530 mAh·g−1 at 0.2 A·g−1, with excellent high-rate performance.
Ball milling is a foundational method for preparing silicon–carbon composites and can achieve uniform mixing. However, it is primarily a physical composite process and cannot effectively suppress silicon volume expansion or construct stable interfaces. Long-cycle performance is therefore limited. In advanced silicon–carbon anodes for lithium-ion cells, ball milling is often combined with CVD to control cost while improving interfacial stability and electrochemical performance. Table 2 summarizes the main silicon–carbon composite preparation methods.
| Method | Carbon uniformity | Interface quality | Cost | Scalability | Main limitation |
|---|---|---|---|---|---|
| Chemical vapor deposition | High | Excellent | High | Low | Complex equipment and slow deposition |
| Liquid impregnation and pyrolysis | Moderate to high | Moderate | Moderate | High | Carbon layer uniformity and interfacial stability |
| Ball milling | Moderate | Poor | Low | High | Physical contact and limited long-cycle stability |
2. Structural Design of Porous Silicon–Carbon Composites
The electrochemical performance of porous silicon–carbon composites is critically determined by their microscopic architecture. Spatial configuration directly governs charge transport, stress distribution, and interfacial stability. In this section, I analyze three mainstream structural families: core–shell, yolk–shell, and embedded/encapsulated structures. I also discuss how structural innovations synergistically address volume expansion, electronic conductivity, and interfacial stability in lithium-ion cells.
The mechanical stress generated during lithiation can be expressed using a simplified core–shell model. For a spherical silicon core with radius \(R\) and a carbon shell with thickness \(h\), the hoop stress in the shell can be approximated as:
$$ \sigma_{\theta} = \frac{E_c}{1-\nu_c} \frac{\Delta R}{R} $$
where \(E_c\) is the Young’s modulus of the carbon shell, \(\nu_c\) is Poisson’s ratio, and \(\Delta R/R\) is the radial strain. When \(\sigma_{\theta}\) exceeds the fracture strength of the carbon shell, cracking occurs. This principle underlies the design of void space in yolk–shell structures and the use of flexible graphene in core–shell structures.
2.1 Core–Shell Structure
A core–shell structure uses a continuous carbon layer to fully coat a silicon core. The carbon shell builds an efficient electron transport path and confines solid electrolyte interphase formation to the carbon surface. The central challenge is the mechanical mismatch between the large volume expansion stress of silicon and the rigidity of the carbon shell, which can cause fatigue cracking. Depending on the dimensionality of the carbon shell, core–shell structures are divided into zero-dimensional amorphous carbon coating and two-dimensional graphene flexible coating.
2.1.1 Zero-Dimensional Amorphous Carbon Coating
Dense amorphous carbon shells are a classic choice for uniform coating of silicon particles, forming a typical core–shell structure. The carbon layer physically isolates silicon from the electrolyte, suppresses side reactions and continuous solid electrolyte interphase growth, improves initial Coulombic efficiency and cycle life, and provides continuous electron transport pathways. It can also partially buffer the mechanical stress of silicon lithiation.
For example, a zero-dimensional core–shell structure \(\mathrm{Ag-PSi@SiC@C}\) was constructed through silver nanoparticle modification, glucose pyrolysis, and high-temperature treatment. This anode retained a specific capacity of 962.6 mAh·g−1 after 200 cycles at 2.0 A·g−1. Beyond interface reinforcement, optimizing the intrinsic properties of the amorphous carbon shell can also significantly improve electrochemical performance. A polyacrylonitrile-derived nitrogen-doped carbon layer was used to coat micrometer-sized porous silicon, forming \(\mathrm{P-Si@C-PAN}\). This composite delivered 685 mAh·g−1 at 8 A·g−1. The structure–performance relationship is based on the dual role of rigid encapsulation and interfacial isolation.
However, the huge volume expansion of the silicon core during lithiation imposes periodic expansion stress on the rigid carbon shell. When the stress exceeds the fatigue strength of the carbon shell, microcracks or local fracture occur. This can trigger a chain reaction: structural damage → electrolyte penetration → silicon core exposure → repeated solid electrolyte interphase rupture and regeneration → continuous consumption of active lithium and electrolyte → rapid capacity fade in lithium-ion cells. To overcome these limitations, I identify several improvement strategies: constructing composite or gradient carbon shells with graphene or carbon nanotubes; designing mesopores or reserved internal buffer space in the carbon shell; and enhancing the binding strength between the carbon shell and silicon core through chemical bonding or interface engineering.
2.1.2 Two-Dimensional Graphene Flexible Coating
Compared with rigid amorphous carbon shells, two-dimensional graphene materials possess excellent flexibility and high electronic conductivity. They can provide flexible encapsulation of silicon particles, effectively accommodating volume expansion during charge and discharge. The high electron mobility of graphene also creates fast in-plane conductive channels, compensating for the low conductivity of silicon.
Graphene is therefore an ideal flexible coating for anodes in lithium-ion cells. For example, a graphene-wrapped porous silicon composite, \(\mathrm{mpSi@G}\), was prepared by magnesiothermic reduction, spray drying, and high-temperature calcination. It delivered a reversible capacity of 1109.5 mAh·g−1 at 5.0 A·g−1. Because a single carbon shell has limitations in mechanical and interfacial performance, adding a second functional shell outside the graphene flexible encapsulation is an effective design strategy. A double core–shell composite, \(\mathrm{P-Si@G@sc-TiO_2}\), was fabricated by electrostatic self-assembly followed by sol–gel deposition of an ultrathin \(\mathrm{TiO_2}\) layer. The inner graphene layer improved conductivity, while the outer \(\mathrm{TiO_2}\) layer enhanced mechanical strength. The material retained 68% of its capacity after 300 cycles at 2 A·g−1.
Unlike the rigid encapsulation of zero-dimensional amorphous carbon, the core mechanism of two-dimensional graphene flexible coating is flexible encapsulation and in-plane conduction. However, graphene sheets easily stack and agglomerate, leading to incomplete coating or multilayer stacking. This not only hinders uniform electrolyte wetting and lithium-ion transport but may also expose the silicon core due to relative sliding between sheets. The failure chain can be described as graphene stacking → blocked ion/electron transport → interfacial delamination → silicon core exposure → continuous solid electrolyte interphase growth → capacity fade in lithium-ion cells.
To overcome these limitations, several approaches are used: chemical bonding to enhance graphene–silicon interfacial adhesion; introducing carbon nanotubes as spacers between graphene sheets to prevent stacking and ensure complete electron transport paths; and introducing hydrophilic functional groups on graphene to improve electrolyte wetting and uniform lithium-ion transport. Table 3 compares representative core–shell structures.
| Structure | Preparation route | Initial Coulombic efficiency (%) | Discharge capacity (mAh·g−1) | Cycle number | Current density (mA·g−1) | Capacity retention (%) |
|---|---|---|---|---|---|---|
| \(\mathrm{Ag-PSi@SiC@C}\) | Metal modification + pyrolysis + high-temperature treatment | 90.10 | 962.6 | 200 | 2000 | 26.5 |
| \(\mathrm{P-Si@C-PAN-400}\) | Solid-state reaction + wet chemistry + heat treatment | 87.35 | 857.6 | 200 | 4000 | 67.5 |
| \(\mathrm{mpSi@G}\) | Magnesiothermic reduction + spray drying + calcination | 70.20 | 2974.9 | 100 | 1000 | 66.0 |
| \(\mathrm{P-Si@G@sc-TiO_2}\) | Electrostatic self-assembly | 80.59 | 1005.1 | 300 | 2000 | 68.0 |
2.2 Yolk–Shell Structure
A yolk–shell structure creates a controlled void between the silicon core and the carbon shell, changing the stress management mode. The void acts as a buffer space for silicon expansion, effectively relieving mechanical stress and protecting the outer carbon shell from fracture. The related structural designs are mainly divided into two categories: amorphous carbon shell with void space, and graphene-based flexible yolk–shell structures.
2.2.1 Amorphous Carbon Shell with Void Space
To alleviate the volume expansion of silicon anodes during charge and discharge, reserving void space between the silicon core and the amorphous carbon shell is an effective strategy. The carbon shell provides a conductive path and constrains silicon expansion, while the void provides buffer space for volume change, avoiding shell rupture and material pulverization. This significantly improves cycling stability in lithium-ion cells.
For example, a yolk–shell structure \(\mathrm{pSi@void@NMC}\) was prepared from an MFI-type pure-silicon zeolite precursor through magnesiothermic reduction, silica coating, and etching. Finite element analysis showed lower lithiation stress and strain. The electrode exhibited a thickness expansion of only 14.6% after 300 cycles. In another approach, porous silicon nanospheres were prepared by sodium thermal reduction, coated with an \(\mathrm{Al_2O_3}\) sacrificial layer by atomic layer deposition, and then encapsulated by electrospun carbon fibers. After acid etching, a 40 nm void was formed, producing a yolk–shell \(\mathrm{pSi@void@CNFs}\). This material retained 952 mAh·g−1 after 200 cycles at 0.2 A·g−1.
Unlike the rigid carbon shell directly coating the silicon core in core–shell structures, the yolk–shell structure reserves a controllable void between them, allowing silicon expansion to be released into the internal void and preventing the carbon shell from bearing periodic stress. However, the void is an inactive region, which reduces the tap density and volumetric energy density of the electrode:
$$ \rho_{tap} = \frac{m}{V_{tap}} $$
$$ E_v = \frac{Q_{rev} V_{avg}}{V_{electrode}} $$
Moreover, the contact between the silicon core and carbon shell changes from areal contact to point contact. If the void is too large, some silicon cores may lose effective electrical contact. In addition, the carbon shell may still crack due to electrolyte corrosion or stress fatigue during long-term cycling. Once broken, the electrolyte can directly attack the silicon core, causing continuous solid electrolyte interphase formation and eventual electrochemical performance decay in lithium-ion cells.
To overcome these shortcomings, I identify the following strategies: optimizing void design through combined theoretical calculation and experiment to determine the best void fraction and distribution; constructing carbon nanotube/graphene-reinforced composite carbon shells to improve mechanical toughness and ionic conductivity; and preventing silicon–carbon delamination through surface modification or simultaneous carbonization and introduction of strong chemical bonds.
2.2.2 Graphene-Based Flexible Yolk–Shell Structure
Traditional yolk–shell silicon–carbon materials suffer from poor conductivity and insufficient mechanical strength because of limited contact between the silicon core and carbon shell. To address this, researchers introduce graphene layers between them. Graphene acts as a flexible conductive skeleton to accelerate electron transport and relieve silicon volume stress. The internal void accommodates silicon expansion, while the outer carbon shell ensures electrode interfacial stability.
For example, vertical graphene sheets were grown on silicon nanospheres by CVD to construct a multidimensional conductive yolk–shell structure. A battery prepared with this material achieved a volumetric energy density of 1083 Wh·L−1. In another design, a ZIF-67-derived carbon shell was used to construct \(\mathrm{Si@C-ZIF}\) yolk–shell units. After heat treatment to grow carbon nanotubes and assemble reduced graphene oxide, a \(\mathrm{Si@C-ZIF@CNTs@rGO}\) composite was obtained. The internal shell buffered expansion and isolated the electrolyte, while the external structure formed a three-dimensional conductive network. This material retained a specific capacity of 1925.2 mAh·g−1 after 300 cycles at 1 A·g−1.
Unlike traditional rigid carbon-shell yolk–shell structures, graphene-based structures use high flexibility and conductivity to build an intermediate layer that combines stress buffering and charge transport. However, several drawbacks remain: graphene sheets easily stack, interrupting conductive paths; physical adsorption with the silicon core leads to easy delamination during long-term cycling; and hydrophobicity hinders electrolyte wetting, affecting lithium-ion transport. The failure chain may be graphene stacking → local conductive network interruption → interfacial delamination → silicon core exposure → continuous solid electrolyte interphase growth → capacity and rate performance decay in lithium-ion cells.
To overcome these drawbacks, the main improvement methods include constructing Si–C or Si–O–C covalent bonding layers between graphene and silicon; compositing graphene with amorphous carbon and carbon nanotubes to form multilayer shells with both flexibility and density; and optimizing the yolk–shell spacing through theoretical calculations to balance expansion space and energy density. Table 4 summarizes representative yolk–shell structures.
| Structure | Preparation route | Initial Coulombic efficiency (%) | Discharge capacity (mAh·g−1) | Cycle number | Current density (mA·g−1) | Capacity retention (%) |
|---|---|---|---|---|---|---|
| \(\mathrm{pSi@void@NMC}\) | Magnesiothermic reduction + silica coating + etching | 77.80 | 1769.8 | 300 | 200 | 95.1 |
| \(\mathrm{pSi@void@CNFs}\) | Atomic layer deposition + electrospinning + carbonization | 71.00 | 952.0 | 200 | 200 | 48.0 |
| \(\mathrm{VGSs-YS-Si/C}\) | CVD + multistep carbonization + polymer self-assembly | 88.30 | 1683.2 at 0.1 C | 1000 at 1 C | — | 80.1 |
| \(\mathrm{Si@C-ZIF@CNTs@rGO}\) | In situ catalytic heat treatment + electrostatic self-assembly | 67.66 | 1925.2 | 300 | 1000 | — |
2.3 Embedded/Encapsulated Structure
The design goal of embedded or encapsulated structures is to fix porous silicon nanoparticles inside a continuous conductive carbon matrix, thereby constructing a stable conductive and buffering skeleton. Typical designs include one-dimensional carbon nanotube network structures, two-dimensional graphene interlayer embedded/encapsulated structures, and three-dimensional porous carbon skeleton supported structures.
2.3.1 One-Dimensional Carbon Nanotube Network
Core–shell and yolk–shell structures use carbon as an outer shell to wrap silicon cores. In contrast, one-dimensional carbon network embedded structures encapsulate multiple silicon nanoparticles within a single carbon nanotube or carbon nanofiber. This structure provides efficient directional electron transport channels and relies on the carbon wall or fiber layer for mechanical confinement. It buffers silicon volume expansion, maintains stable electrical contact, and improves electrode structural stability and cycle life in lithium-ion cells.
For example, layered porous silicon microparticles were prepared by magnesiothermic reduction and coated with carboxylated carbon nanotubes through electrostatic adsorption, forming a conductive carbon nanotube cage. This material retained 91.9% of its capacity after 500 cycles at 1 A·g−1. To further improve electrode mechanical integrity and build a more efficient conductive network, a multilevel encapsulation strategy was used. Porous silicon was electrostatically self-assembled with graphene oxide and carbon nanotubes to form a three-dimensional interpenetrating framework. The resulting \(\mathrm{pSi@GO-CNTs@C}\) anode retained 221 mAh·g−1 after 1000 cycles at 5 A·g−1.
Overall, this structure maintains good electrical contact and reduces contact resistance compared with core–shell and yolk–shell structures. However, carbon nanotubes easily agglomerate during composite preparation, leading to nonuniform network distribution. In addition, the nanotube network may undergo fatigue fracture or structural collapse under repeated stress, especially at high silicon loading, where local stress concentration accelerates network destruction. The failure chain can be described as nanotube agglomeration or nonuniform network → local conductive interruption → interfacial delamination → capacity fade in lithium-ion cells.
Current improvement methods focus on multidimensional composites, structural design, and interface strengthening: constructing multidimensional hybrid conductive networks, such as using two-dimensional sheets to prevent carbon nanotube agglomeration and enhance mechanical integrity; developing scalable uniform dispersion processes so that carbon nanotubes interweave uniformly inside and outside microspheres; and strengthening interfacial bonding through carbon nanotube surface functionalization or introducing a SiC transition layer.
2.3.2 Two-Dimensional Graphene Interlayer Embedded/Encapsulated Structure
Flexible encapsulation technology is a key means to solve the volume expansion problem of silicon-based materials during cycling. By constructing a three-dimensional continuous graphene network, the active material is encapsulated and constrained as a whole. This enables fast electron conduction and uses graphene sheet sliding and deformation to release internal stress, ensuring structural stability during charge and discharge in lithium-ion cells.
For example, a nano–micro composite electrode structure was designed in which a lithium phosphate artificial interface layer and a carbon layer were sequentially constructed on porous silicon; the outer layer combined graphitic carbon and polyethylene oxide to build ion channels and suppress electrode expansion. This anode retained 941.5 mAh·g−1 after 100 cycles. In addition, using the flexibility of low-dimensional carbon materials for physical stress management can also alleviate silicon-based expansion. A \(\mathrm{p-SiO@rGO}\) composite anode was designed with finite element simulation guidance. This material retained 589.7 mAh·g−1 after 1000 cycles.
Compared with the line contact of one-dimensional carbon nanotube networks, two-dimensional graphene sheets provide a more continuous and complete conductive network. However, the intrinsic contradiction of this structure is that graphene sheets easily stack and agglomerate during composite preparation, reducing interlayer spacing and weakening stress buffering and ion transport. In addition, graphene and silicon particles are mostly physically adsorbed, lacking strong chemical bonding, so the interface easily delaminates during long-term cycling. The failure chain may be graphene stacking or agglomeration → compressed interlayer spacing → blocked ion transport → interfacial delamination → silicon particle exposure → continuous solid electrolyte interphase growth → capacity fade in lithium-ion cells.
To improve these shortcomings, the following measures can be taken: interfacial strengthening by constructing Si–O–C covalent bonds or introducing functional intermediate layers such as \(\mathrm{Li_3PO_4}\) to convert physical adsorption into strong chemical bonding; structural optimization by constructing multilevel composite networks to prevent graphene stacking and form a three-dimensional continuous conductive skeleton; and theoretical guidance using finite element simulation to predict and design key parameters such as interlayer spacing.
2.3.3 Three-Dimensional Porous Carbon Skeleton Supported Structure
Unlike the previous two approaches, the three-dimensional porous carbon skeleton supported structure does not use carbon to coat silicon. Instead, it disperses and anchors porous silicon in a three-dimensional interconnected carbon skeleton. This structure forms a continuous conductive network and uses high mechanical strength and porosity to share silicon expansion stress, separate silicon particles, inhibit agglomeration and stress superposition, and significantly improve structural stability in lithium-ion cells.
For example, a zeolitic imidazolate framework-8 layer was grown on silicon nanoparticles and carbonized to form a porous carbon layer. Then, porous silicon, graphene oxide, and sucrose were assembled by solvothermal treatment and heat treatment to prepare a graphene-supported double-layer carbon-encapsulated porous silicon composite. The capacity retention reached 89.5% after 240 cycles. In another route, a melt-processable polyacrylonitrile copolymer was used as a carbon precursor. After melt extrusion and microwave irradiation, a \(\mathrm{p-Si@C}\) composite was prepared, and Si–C bonds formed at high temperature to improve interfacial bonding. When mixed with graphite at a 1:9 ratio, the material retained a reversible capacity of 535.1 mAh·g−1 after 300 cycles.
However, the intrinsic contradiction of this structure is that silicon loading is limited. If the silicon content is too high, the skeleton may rupture or the pore channels may become blocked. In addition, interfacial bonding is weak because silicon particles are mostly physically embedded in the carbon skeleton without strong chemical bonding, so they may loosen and detach during long-term cycling. At the same time, the carbon skeleton is at risk of fatigue collapse under repeated stress, especially when silicon loading is nonuniform, where local stress concentration accelerates skeleton destruction. The failure chain can be described as nonuniform or excessive silicon loading → pore channel blockage or skeleton rupture → interfacial loosening → silicon particle detachment → capacity fade in lithium-ion cells.
To address these shortcomings, recent research has formed a clear optimization path: first-generation optimization of the carbon skeleton structure by tuning pore size to anchor silicon particles; second-generation strengthening of interfacial bonding through covalent bonds for chemical anchoring; and third-generation construction of multidimensional network synergy by compositing carbon materials of different dimensions to form multilevel conductive reinforcements. Table 5 compares representative embedded/encapsulated structures.
| Structure | Preparation route | Initial Coulombic efficiency (%) | Discharge capacity (mAh·g−1) | Cycle number | Current density (mA·g−1) | Capacity retention (%) |
|---|---|---|---|---|---|---|
| \(\mathrm{LSP-Si@CNT}\) | Magnesiothermic reduction + electrostatic adsorption | 84.40 | 1862.1 at 100 mA·g−1 | 500 | 1000 | 91.9 |
| \(\mathrm{pSi@GO-CNTs@C}\) | Electrostatic self-assembly + CVD | 57.30 | 221.0 | 1000 | 5000 | — |
| \(\mathrm{P-MSi+GC7@PEO}\) | Silver-catalyzed chemical etching + heat treatment + carbonization | 80.50 | 941.5 | 100 | 1000 | 32.0 |
| \(\mathrm{p-SiO@rGO}\) | Finite element simulation guided | 76.40 | 589.7 | 1000 | 2000 | 91.4 |
| \(\mathrm{Graphene/IOC@pSi}\) | Ball milling + solvothermal reaction + carbonization | 60.60 | 1182.0 | 240 | 200 | 89.5 |
| \(\mathrm{p-Si@C-G}\) | Melt extrusion with micrometer silicon + microwave irradiation | 75.60 | 535.1 | 100 | 268 | 84.1 |
2.4 Failure Mechanisms of Porous Silicon–Carbon Anodes
Although core–shell, yolk–shell, and embedded/encapsulated structures each have distinctive features in mitigating silicon volume expansion, building conductive networks, and stabilizing electrode interfaces, they still face different failure pathways during long-term cycling in lithium-ion cells. Clarifying these failure mechanisms is a prerequisite for further optimizing structural design and improving cycle life. Table 6 compares the main failure modes of the three structure types.
| Structure type | Main failure mode | Failure cause | Electrochemical consequence |
|---|---|---|---|
| Core–shell | Carbon shell fatigue fracture and interfacial delamination | Silicon expansion stress exceeds carbon shell fatigue limit | Continuous SEI thickening and rapid capacity fade |
| Yolk–shell | Loss of electrical contact and local carbon shell cracking | Unstable electrical contact due to void and carbon shell defects | Rate performance decline and sudden capacity drop in later cycles |
| Embedded/encapsulated | Interfacial peeling and carbon skeleton collapse | Weak interfacial bonding and nonuniform silicon loading | Decreased cycling stability and overall electrode failure |
2.4.1 Failure Mechanism of Core–Shell Structures
Core–shell structures use a dense carbon layer for rigid encapsulation of the silicon core, confining the unstable solid electrolyte interphase to the outer carbon surface. However, the huge volume expansion of silicon during lithiation imposes periodic expansion stress on the carbon shell. When the stress exceeds the fatigue strength of the carbon shell, microcracks or local fracture occur. The electrolyte penetrates through the cracks and directly contacts the silicon core, causing repeated rupture and regeneration of the solid electrolyte interphase on the silicon core surface. This process continuously consumes active lithium and electrolyte and also leads to delamination between the carbon shell and silicon core. The electron transport path is interrupted, and the final result is rapid capacity fade in lithium-ion cells.
The growth of the solid electrolyte interphase can be approximated by a parabolic law:
$$ t_{SEI} \propto \sqrt{D_{SEI} t} $$
where \(t_{SEI}\) is the solid electrolyte interphase thickness, \(D_{SEI}\) is the diffusion coefficient of reactive species through the interphase, and \(t\) is time. Repeated fracture exposes fresh silicon, resetting \(t_{SEI}\) and accelerating electrolyte consumption. The interfacial resistance can be written as:
$$ R_{int} = R_{ct} + R_{SEI} $$
where \(R_{ct}\) is the charge-transfer resistance and \(R_{SEI}\) is the solid electrolyte interphase resistance. Both increase as the interphase thickens and delamination proceeds, degrading the rate capability of lithium-ion cells.
2.4.2 Failure Mechanism of Yolk–Shell Structures
Yolk–shell structures reserve a void between the silicon core and carbon shell, effectively dissipating volume expansion stress into the buffer space and preventing the carbon shell from directly bearing tensile stress. However, two typical failure modes still exist during long-term cycling. First, the presence of voids reduces the tap density of the electrode. If the void fraction is too large, point contacts or short-range conductive bridges between the carbon shell and silicon core may break due to relative displacement during cycling, causing some silicon cores to lose electrical contact. Second, although the carbon shell does not directly bear expansion stress, it may still undergo fatigue cracking due to electrolyte corrosion or local stress concentration. This is especially true when the carbon shell thickness is nonuniform or structural defects exist. The solid electrolyte interphase may then continue to grow on the inner surface of the carbon shell, blocking lithium-ion transport channels in lithium-ion cells.
The void fraction can be defined as:
$$ f_v = \frac{V_{void}}{V_{core} + V_{void} + V_{shell}} $$
An optimal \(f_v\) must balance expansion accommodation and volumetric energy density. The capacity retention after \(N\) cycles is commonly expressed as:
$$ R_{ret} = \frac{Q_N}{Q_1} \times 100\% $$
where \(Q_N\) is the discharge capacity at cycle \(N\) and \(Q_1\) is the initial discharge capacity. For yolk–shell structures with poorly designed voids, \(R_{ret}\) may drop sharply after several hundred cycles because of electrical disconnection and localized solid electrolyte interphase growth.
2.4.3 Failure Mechanism of Embedded/Encapsulated Structures
Embedded/encapsulated structures disperse porous silicon particles in a continuous carbon nanotube network, graphene sheets, or porous carbon matrix. The stress is borne by the skeleton as a whole, and electron transport paths are abundant. The main failure mechanisms include the following. First, interfacial delamination between silicon particles and the carbon skeleton occurs because the thermal expansion coefficients of silicon and carbon differ significantly. If the interface relies only on physical adsorption without chemical bonding, interfacial peeling easily occurs during cycling, causing silicon particles to detach from the conductive network. Second, the carbon skeleton may undergo structural collapse or local fracture under repeated stress. This is especially true when silicon loading is too high or nonuniformly distributed, where local stress concentration accelerates skeleton destruction and leads to overall electrode structural instability in lithium-ion cells.
The capacity fade rate can be approximated by an empirical relation:
$$ Q(N) = Q_0 \exp(-\lambda N) $$
where \(Q_0\) is the initial capacity, \(\lambda\) is a fade constant related to mechanical and interfacial degradation, and \(N\) is the cycle number. Minimizing \(\lambda\) requires strong interfacial bonding, uniform silicon dispersion, and a robust carbon skeleton. Table 7 summarizes the structure–performance–failure relationships and the corresponding design strategies.
| Structure | Stress management | Conductive network | Main weakness | Design strategy |
|---|---|---|---|---|
| Core–shell | Rigid carbon shell constrains expansion | Continuous carbon coating | Carbon shell fatigue and delamination | Gradient shell, chemical bonding, buffer pores |
| Yolk–shell | Internal void buffers expansion | Carbon shell plus internal contact | Low tap density and contact loss | Optimized void fraction, reinforced shell |
| Embedded/encapsulated | Carbon skeleton shares stress | 1D, 2D, or 3D carbon network | Interfacial peeling and skeleton collapse | Covalent bonding, uniform loading, multilevel network |
3. Summary and Outlook
In this review, I have analyzed porous silicon–carbon composite anodes for lithium-ion cells from the perspectives of preparation chemistry and spatial structure design. Based on structural design principles and stress regulation mechanisms, I classify these materials into core–shell structures, yolk–shell structures, and embedded/encapsulated structures. According to the dimensionality of the carbon matrix, they can be further divided as follows: core–shell structures include zero-dimensional amorphous carbon dense coating and two-dimensional graphene flexible encapsulation; yolk–shell structures include zero-dimensional amorphous carbon shells with void space and two-dimensional graphene-based flexible conductive layers; and embedded/encapsulated structures include one-dimensional carbon nanotube line-contact networks, two-dimensional graphene interlayer confinement, and three-dimensional porous carbon skeletons for overall load bearing.
The three structure families have different application scenarios in porous silicon–carbon composite anodes for lithium-ion cells. Core–shell structures achieve high volumetric energy density through dense carbon layers and are suitable for high-tap-density systems, but rigid carbon shells are prone to fatigue failure. Yolk–shell structures dissipate expansion stress through reserved voids, maintain carbon shell integrity and stable solid electrolyte interphase, and exhibit excellent long-cycle life, making them suitable for high-stability energy storage systems. However, they sacrifice some volumetric energy density and still face risks of conductive path failure. Embedded/encapsulated structures disperse silicon particles in a three-dimensional carbon skeleton, providing abundant electron transport paths and outstanding rate performance, making them suitable for high-rate scenarios. Nevertheless, their interfacial stability depends strongly on the silicon–carbon bonding strength.
For future development, I believe that structural design of porous silicon–carbon composite anodes should continue to break through in the following directions. First, introduce Si–C, Si–O–C, and other covalent bonds to replace physical contact and strengthen interfacial bonding. Second, integrate core–shell, yolk–shell, and embedded/encapsulated structures to construct multilevel composite architectures that simultaneously achieve high volumetric energy density, long cycle life, and high rate performance in lithium-ion cells. Third, develop green synthesis routes using industrial silicon waste as raw material and avoiding expensive templates, combined with continuous processing, to promote low-cost scalable manufacturing. In addition, the success of current experiments is sometimes serendipitous. Therefore, promoting deep integration of theoretical calculations and experiments is an effective way to improve research efficiency and success rate in the development of advanced lithium-ion cells.
In summary, porous silicon–carbon composites represent one of the most promising pathways toward next-generation high-energy lithium-ion cells. The key is to balance porosity, carbon dimensionality, interfacial chemistry, and mechanical integrity. When these factors are optimized synergistically, silicon-based anodes can move closer to practical deployment in electric vehicles, portable electronics, and grid-scale energy storage. Table 8 outlines the major future research directions and their expected impact on lithium-ion cell performance.
| Future direction | Specific action | Expected benefit for lithium-ion cells |
|---|---|---|
| Interfacial covalent bonding | Introduce Si–C and Si–O–C bonds | Suppress delamination and improve cycle life |
| Multilevel composite architecture | Combine core–shell, yolk–shell, and embedded structures | Balance energy density, rate capability, and stability |
| Green and scalable synthesis | Use industrial silicon waste and template-free routes | Reduce cost and enable mass production |
| Theory–experiment integration | Use finite element simulation and machine learning | Accelerate optimization of pore size, void fraction, and interface |
| Advanced characterization | In situ TEM, XRD, and neutron depth profiling | Reveal real-time degradation mechanisms |
Through these efforts, I expect that porous silicon–carbon anodes will gradually overcome their current limitations and become a key enabling technology for high-performance lithium-ion cells. The transition from laboratory-scale innovation to industrial-scale manufacturing will require not only materials breakthroughs but also electrode engineering, electrolyte optimization, and cell-level design. Nevertheless, the combination of porous silicon and carbon matrices remains one of the most rational and versatile strategies for harnessing the extremely high capacity of silicon while maintaining the reliability expected from modern lithium-ion cells.