Recent advances and perspectives of microsized alloying-type porous anode materials in high-performance Li- and Na-ion batteries

Dealloying methods

Dealloying is often expressed as a bimetallic or multi-metal alloy corrosion process, in which more active metals in the precursors react selectively with the medium and are removed operando in the following steps to leave a more stable component. During the reaction, the residual component diffuses between the parent alloy and reactive medium to form a bicontinuous porous structure with 3D open pores and maximum interfaces^[34]. Dealloying has attracted intensive attention for the fabrication of porous alloys because pore creation does not require additives and adjustable structures. The 3D porous alloying metal supplies rich Li and Na storage and active sites and enables rapid mass transfer, which is vital for LIBs and SIBs. According to different alloy precursors, the porosity of the structure can be adjusted from 40% to 80% by taking advantage of the alloy parting limit^[34,35]. The dealloying of porous structures is carried out by chemical, electrochemical, low-melting metal, and other dealloying methods.

Chemical dealloying is one of the most common methods to obtain porous structures from the alloy precursor via spontaneous chemical reactions based on the different standard electrode potentials in an acidic or alkaline solution. The precursor alloy must be homogeneous or have an intermetallic phase comprising a relatively active metal and another more inert one. Porous metals containing Si, Ge, Sn, Sb, and so on have been prepared by dealloying routes. For example, Feng et al. fabricated porous Si consisting of nanodendrites by direct dealloying of a commercial Al-Si eutectic alloy with microparticles in 5 M HCl for 24 h^[36]. Zhang et al. prepared porous microscale antimony from Mg-Sb alloys by one-step chemical vapor dealloying with a HCl vapor as the dissolution medium at 600 °C [Figure 5A]^[37]. Sohn et al. prepared porous Si via the dealloying of Si/Al-Cu-Fe nanoparticles obtained from smelting, roller melt-spinning, and attrition milling via dual-step acidic and alkaline wet chemical etching^[38]. Lin et al. reported a low-temperature molten salt system for the metallothermic (Mg or Zn) reduction of GeO₂ to obtain mesoporous hollow Ge microspheres, and the molten ZnCl₂ promotes the reaction at a low temperature. In the etching process, part of the metal alloy matrix is consumed due to acid leaching to generate abundant pores, and subsequent alkaline etching decreases the Si domain size and enlarges the pores^[39].

Figure 5. Schematic illustrations of porous metal preparation via dealloying fabrication methods. (A) Schematic illustration of fabrication procedures of P-Sb. Reproduced with permission^[37] (Copyright 2020, Elsevier B.V.). (B) Preparation of bulk 3D porous-Si by dealloying in a metallic melt. Reproduced with permission^[47] (Copyright 2014, ACS Publications). (C) Schematic illustration of fabrication procedures of 3D porous Si. Reproduced with permission^[48] (Copyright 2016, Elsevier B.V.). (D) Schematic illustration for the multi-step synthesis of ultrafine Si porous structures. Reproduced with permission^[49] (Copyright 2020, Kluwer Academic Publishers).

In addition to single metals, porous bimetallic materials have also been obtained by chemical dealloying using ternary alloys as the raw materials. Yang et al. produced Si-Ge bimetallic porous materials via the dealloying of Al-Ge-Si microparticles with an Al content of 80% in HCl^[40]. Other bimetallic compounds, such as Bi-Sn^[41], Ag-Ge^[42], SnSb^[43], and SnNi^[44], have been formed by similar dealloying methods. Although chemical dealloying using acidic and alkaline media as the dissolving agents is an emerging route, there are some shortcomings, such as its time-consuming nature, the production of waste acidic or alkaline solutions, and structural collapse of the porous architecture. To solve these problems, electrochemical dealloying methods have been developed to obtain porous Si, Ge, and Ge-Si alloys by adjusting the applied potential. Yin et al. conducted electrochemical dealloying in a molten CaCl₂-NaCl eutectic electrolyte at 600 °C to produce porous Ge^[45]. Yuan et al. adopted Mg alloying/dealloying using a molten MgCl₂-NaCl-KCl electrolyte at 500 °C to fabricate porous Si^[46]. Although the porous structure can be retained, the conditions in electrochemical dealloying are relatively harsh and an oxygen-free environment is needed to prevent oxidation of the porous materials.

Liquid metal dealloying has been proposed to prepare porous alloys because of the controllable porosity. Wada et al. reported the fabrication of bulk 3D porous Si via a metallic melting dealloying route with Mg₂Si as the precursor and Bi metal as the dealloying melting medium^[47]. Mg dissolves selectively from the Si-Mg alloy into the Bi melt to produce porous Si at high temperatures [Figure 5B]. Similarly, Wada et al. fabricated porous 3D Si by immersing a Mg-Si alloy precursor in a Bi melt and washing it with 1 M nitric acid [Figure 5C]^[48]. However, electrochemical and melting metal dealloying have some disadvantages, including environmental issues and economic efficiency, which hinder their wider application.

Alternative dealloying methods have been developed to obtain more desirable porous structures efficiently. For instance, Chen et al. and Wang et al. synthesized porous Si structures with various particle sizes via the oxidation of ball-milled Mg₂Si in Ar/O₂ or air in conjunction with pickling reactions in HCl to remove the MgO impurity [Figure 5D]^[49,50]. We recently developed a new strategy to fabricate ant-nest-like nanoporous Si microparticles by nitriding Mg₂Si particles under nitrogen at 750 °C to generate a Mg₃N₂/Si composite, followed by the dissolution of Mg₃N₂ in HCl. The ant-nest-like porous Si comprises 3D-interconnected Si nano-ligaments and bicontinuous nanopores^[51]. More recently, An et al. fabricated uniform porous Si particles by treating a commercial Mg-Si alloy in CO₂ at 500-700 °C and etching MgO in HCl^[52]. In this process, CO₂ is reduced into the carbon layer, which is in situ coated on the surface of porous Si. The gas dealloying method involving a reactive gas, such as O₂, N₂, or CO₂, reacting with the Si-Mg alloy produces phase separation and generates the 3D composite comprising Si nano-ligaments and the surrounding intermediate product. The porous 3D structure is obtained after dissolving the intermediate product. Compared to the chemical/electrochemical dealloying method, the gas dealloying technique not only can avoid structural collapse during etching but is also more environmentally friendly and faster.

Template method

The template method is one of the most common methods to construct porous materials because the product has the inverse structure of the porous template. Moreover, the size of the porous structure can be adjusted by controlling the template parameters. The common templates are opals, organic spheres, and SiO₂. The template method proceeds by the following steps. Infiltration of the precursor by chemical vapor deposition (CVD) or the sol-gel method is performed, followed by elimination of the matrix either by thermal annealing or chemical etching^[53]. For example, Song et al. used silica opal as a porous template to prepare Ge inverse opal with porous walls by CVD and a Ge wall layer on the silica opal^[54]. The silica template was then etched selectively with HF. Geier et al. produced porous Ge structures by etching and annealing the Ge precursor-loaded poly(methyl methacrylate) template^[55]. The template was eliminated by heating at 500 °C for 5 min under a vacuum. In addition to Ge and Si, regular porous structures can be prepared by the deposition of a Si layer on a SiO₂ opal template with Si₂H₆ as the gas precursor, followed by dissolution in HF^[56]. However, in addition to the complicated preparation process and high-cost template, the template method for porous structures often involves highly toxic acids or emits a large amount of CO₂^[57,58].

To overcome these drawbacks, the self-template synthesis of porous materials has been developed using metal oxides, such as SiO₂, GeO₂, and SnO₂, as both the raw materials and self-sacrificial templates^[59,60]. For instance, Bao et al. converted microscale diatomite into porous 3D microscale Si with continuous nano-ligaments at a relatively low temperature of 650 °C by the following reaction^[59]:

2Mg + SiO₂ → 2MgO + Si (1)

Some natural or synthetic porous SiO₂ materials, such as zeolite SSZ-13^[61], have been used in the fabrication of porous Si. The by-product of MgO is usually used as the pore-forming agent, which can generate uniform nanoscale pores after dissolution in diluted HCl^[62]. Similarly, porous Ge can be obtained by magnesiothermic reduction (MgR) with GeO₂ as the self-sacrificial template^[63]. In addition to MgR, aluminothermic reduction in molten salt can convert the SiO₂ template into porous Si at a low temperature below 300 °C. For example, porous crystalline Si has been synthesized by the aluminothermic reduction of an AlCl₃ salt at 250 °C^[64], and other reducing agents, such as Zn and NaBH₄, have been employed to convert GeO₂ into porous Ge^[65].

The soluble salt-template method has been developed to prepare porous alloys because the soluble salts can be easily removed. The pores are produced by removing the salt by-products of the template rather than chemical etching, thus providing both high porosity and controllability. In addition, the surface area and pore size of the materials can be adjusted by selecting appropriate alloy precursors. Some inorganic salts, such as NaCl, KCl, and LiCl, have been utilized as templates^[66]. Dai et al. reported a bottom-up approach for mesoporous crystalline Si with tunable particle/pore size by SiCl₄ reduction [Figure 6A]^[67]. In this process, SiCl₄ is the Si precursor and NaK is the reducing agent. Si and the salt by-products (NaCl and KCl) are considered to precipitate simultaneously to form a composite nano-matrix due to the poor solubility of NaCl/KCl and Si in the solvent. In the subsequent heat treatment, the amorphous Si is transformed into a crystalline framework composed of interconnected primary particles. The product is washed with water to remove the by-products of salts to form porous Si. Tang et al. designed a series of mesoporous Ge materials by the salt-templating method with NaCl and KCl as the in situ templates [Figure 6B]^[68]. The surface area and pore size and volume can be adjusted by changing the reaction time, temperature, and calcination temperature, and the salt by-products as internal self-forming templates can be removed without an etchant. Compared to traditional template methods, this method is simpler, faster, and more environmentally friendly.

Figure 6. (A) Schematic of synthesis route of mesoporous crystalline Si. Reproduced with permission^[67] (Copyright 2014, Nature Publishing Group). (B) Synthetic route of mesoporous Ge. Reproduced with permission^[68] (Copyright 2020, Elsevier Inc.).

Chemical etching method

Chemical etching is an effective method to fabricate porous alloys based on different galvanic displacements or chemical activities^[69]. Etching methods can be mainly divided into electrode and electroless etching. Electrode etching is controlled by the current, whereas the electroless process often involves other metal ions, such as Ag⁺ and Cu^2+[70-72]. For example, porous Si nanowires have been produced with silver nitrate in HF and doped Si^[73,74]. Typically, two reactions take place during etching^[69]:

4Ag⁺ + 4e^- → 4Ag (2)

Si + 6F^- → [SiF₆]^2- + 4e^- (3)

The porosity and pore size of porous Si can be controlled by the current density and concentration of the etchant. Zhong et al. studied the mechanism of the formation of porous Si nanowires in a two-step Ag-assisted electroless chemical etching method [Figure 7A]^[75]. The H-terminated Si was coated with Ag by electroless deposition or physical vapor deposition and then placed into a solution containing HF and H₂O₂ with variable concentrations for different time durations. Zhang et al. synthesized porous Si materials by the ferrite-assisted chemical etching of the Si substrate [Figure 7B]^[76]. In this reaction, ferrite catalyzes the reaction of Si particles with glycol to produce porous Si particles. To avoid the use of the noble metal Ag, some groups have developed a less expensive Cu ion etching method. For example, porous Si has been prepared by the chemical etching of Cu nanoparticles in HF/H₂O₂^[77]. In addition to Si, porous Ge is also easy to fabricate by this method^[78] using an electrolyte of HF:H₂O₂:H₂O (2:3:1). Chemical etching can produce porous Si/Ge with different porosities, but the formation is relatively slow and involves highly corrosive HF.

Figure 7. Schematic illustrations of porous metal preparation via chemical etching method. (A) Schematic illustration of porous Si nanowires. Reproduced with permission^[75] (Copyright 2012, American Chemical Society). (B) Schematic illustration of the formation process of porous Si wafer. Reproduced with permission^[76] (Copyright 2015, Wiley-VCH Verlag).

Electrochemical etching can also produce porous alloys by applying a current^[79]. The starting materials are often wafers, such as Si and Ge, and the electrolyte is typically HF. Uhlirs at Bell Laboratories accidentally found that porous Si could be generated by the electrochemical etching of a Si wafer^[80]. The pore formation mechanism and adjustment of the pore parameters were investigated^[69]. Thakur et al. synthesized a macroporous Si film by electrochemical etching in 48% HF and dimethylformamide using a constant current density of 2 mA cm^-2[81]. The porosity and depth of porous Si can be adjusted by the current density and HF concentration^[82]. Bioud et al. reported the fabrication of porous Ge films by galvanostatic etching using a two-electrode system with a Ge wafer and Pt plate as the working and counter electrodes, respectively^[83]. The electrolyte consisted of HF and anhydrous ethanol. Porous materials obtained by electrochemical etching have high density and purity. Although the electrochemical etching method can produce porous Si and Ge with different porosities using an applied current, it involves the use of highly corrosive HF and can cause environmental concerns^[84].

Self-assembly methods

Self-assembly is attractive in the preparation of porous microstructures with nanostructures, such as nanoparticles, nanowires, and nanosheets. This method is cost-effective, scalable, simple, and can produce particles with a high tap density. Jia et al. designed porous microscale Si spheres composed of nanoscale crystalline Si with a uniform porous structure by self-assembly using the bottom-up microemulsion approach [Figure 8A]^[85]. Jeong et al. adopted a cyclic iodide process to assemble microclusters comprising many kinked Si nanowires with a diameter of less than 20 nm^[86]. The secondary microscale particles had considerable internal porosity, good thermodynamic stability, and high tap density. Jia et al. produced hierarchical porous Si microspheres of 6-8 μm in diameter consisting of Si-coated CNTs by self-assembly [Figure 8B]^[87]. By introducing molten salts, such as KCl and NaCl, a liquid environment was created to absorb heat and increase ion moving ability at a high temperature. Yao et al. prepared mesoporous diatomite-derived Si microspheres with NaCl as the molten liquid interface to grow Si nanocrystals by the self-assembly of mesoporous Si nanocrystals in MgR^[88]. Nevertheless, there is a rate-limiting step in which the conversion of an irregular shape to a nearly spherical shape depends on the concentration of fused salts and liquid magnesium, which may dissolve each other by the interactions of ions and the negative ions in the cavity of these positive ions depend on the vibration frequency of Mg²⁺, Na⁺, and Cl^-[89].

Figure 8. (A) Schematic illustration of preparation process of porous Si. Reproduced with permission^[85] (Copyright 2018, Elsevier B.V.). (B) Schematic showing synthesis of CNT@Si@C microspheres. Reproduced with permission^[87] (Copyright 2020, Nature Publishing Group).

Si-based porous structures for LIBs and SIBs

Si has the unique cubic space group of Fd3m and a large interplanar crystal spacing of 5.472 Å. It is promising anode material for LIBs on account of its high theoretical capacity, abundant resources, low cost, and non-toxicity. Si also has a low delithiation potential of ~0.4 V vs. Li⁺/Li, which is higher than that of graphite (below 0.2 V). In addition, Si reacts electrochemically with Li to form alloys during cycling, thus offering space to increase the Li-ion storage^[90]. Chevrier et al. carried out first principles calculations on the Si-Li system^[91]. Shenoy et al. derived the optimal lattice parameters of the crystalline Li-Si phases using density functional theory (DFT) calculations according to the crystal structures from the Inorganic Crystal Structure Database^[92]. According to different reaction conditions, the phases in the Li-Si system were confirmed as LiSi, Li₁₂Si₇, Li₇Si₃, Li₁₃Si₄, and Li₂₂Si₅. Among them, Li₂₂Si₅ showed the highest total theoretical capacity of 4,200 mAh g^-1[93]. By combining as many as 4.4 Li atoms, Si shows a volume expansion (~300%) and huge stress during lithiation/delithiation, which causes pulverization, unstable SEI layers, structural collapse, and fast capacity fading. Porous microscale Si-based structures can be divided into nano-unit self-assembled porous Si and continuous and connected skeleton porous Si, as shown in Table 3^{[38,46.48,72,88,94-124]}.

Nano-unit self-assembled microporous Si

Si nanomaterials with a small size and high specific surface area also exhibit high capacity and improved cycling stability but suffer from low ICE and tap density. Although it has a high gravimetric capacity of above 2,000 mAh g^-1, the volumetric energy density of nano-Si is even lower than that of commercial graphite. The assembly Si nanostructures into microscale secondary particles can combine the merits of nano- and micromaterials to produce a high tap density and capacity simultaneously^[125]. Moreover, the nano-units self-assembled in microscale porous Si have many pores to accommodate volume expansion and reduce the volume change on the particle level in LIBs^[112]. Porous Si consists of nano-units for lithiation from end-to-end, whereas solid larger Si particles are surface-to-center^[126].

Liu et al. fabricated yolk-shell spherical Si pomegranate microbeads with a diameter of up to 10 μm via the self-assembly of Si@SiO₂ nanoparticles with the aid of emulsifiers [Figure 9A]^[127]. The spherical Si pomegranate microbeads show a high reversible capacity of 2,350 mAh g^-1 at a rate of C/20 (1 C = 4,200 mA g^-1) in conjunction with a capacity retention of 97% after 1,000 cycles at C/2, significantly better than that of Si NPs [Figure 9B]. The enhanced performance of these composites stems from the incorporation of carbon and the smaller particle sizes, with the well-defined internal space allowing Si to expand without changing the secondary particle size. Jia et al. reported porous Si (p-Si) comprised of Si nanoparticles^[85]. The unique porous structure endows the materials with good electrical contact and provides a fast route for the transportation of lithium ions, as well as good internal stress accommodation. As a result, p-Si exhibits a reversible capacity of 2,500 mAh g^-1 at 0.26 A g^-1 with an ICE of 75% and a cycling stability capacity of 1,467 mAh g^-1 at 2.6 A g^-1 after 370 cycles with a capacity retention of 83%. Furthermore, Shi et al. designed a porous microscale structure made of TiO₂-coated Si nanoparticles in building blocks that were further filled with hierarchical Si@TiO₂@C (SA-SiTC)^[128]. In this composite, the carbon filler and TiO₂ layer act as a mechanical buffer and electrolyte reservoir, respectively. The SA-SiTC electrode displayed long cycling stability (842.6 mAh g^-1 after 1,000 cycles at 2 A g^-1), high volumetric capacity (174 mA h cm^-3) and ICE (80.9%), and a stable SEI layer.

Figure 9. (A) Structural evolution of pomegranate Si microparticles before and after electrochemical cycling. (B) Cycling performance of the Si pomegranate and other structures at C/2. Reproduced with permission^[127] (Copyright 2014, Springer Nature). (C) Microscale Si particles incorporating self-kinked Si nanowire clusters. (D) Cycling performance of samples at 0.5 C. Reproduced with permission^[86] (Copyright 2020, Wiley-VCH Verlag). (E) SEM images of Si nanowires and mapping of Si. Reproduced with permission^[129] (Copyright 2020, Elsevier B.V.).

In addition to Si nanoparticles, Si nanowires can be assembled into porous secondary microscale particles. Jia et al. applied self-assembly to prepare hierarchical porous Si microspheres of 6-8 μm in diameter with Si-coated CNTs (CNT@Si@C)^[87]. The CNT@Si@C electrode with a practical mass loading (3 mAh cm^-2) showed a specific capacity of 750 mAh g^-1, an initial swelling of < 20% at 100% state-of-charge and a capacity retention of 92% over 500 cycles. Owing to the unique highly porous yarn-ball-like structure with high porosity and good mechanical strength, the CNT@Si@C microspheres showed the expected effect of suppressed swelling of porous Si particles upon lithiation to maintain the overall structural integrity. Jeong et al. fabricated Si clusters with an average size of 2.7 μm by a silicon iodide (SiI₄) decomposition method, as shown in Figure 9C^[86]. The porous Si was composed of kinked nanowires with diameters of less than 20 nm. The kinked Si nanowires not only formed voids but also contacted each other to yield electrolyte transfer channels for better ionic conductivity and interconnections for better electronic conductivity. The nanowire-assembled porous particles have ample voids between the nanowires and can accommodate large volume changes. As a result, they showed a capacity of 2,784.7 mAh g^-1 at 0.1 C with an ICE of 82.4%, together with a superior capacity retention of 83.6% for 1,000 cycles [Figure 9D]. Recently, Choi et al. produced porous microscale Si consisting of nanowires with a diameter of 300 nm and a length of > 1 μm by self-assembly using a molten salt electrochemical reduction method from RH-derived SiO₂^[129]. As shown in Figure 9E, the product has a spherical shape with a porous structure. The porous Si anode has a high capacity of 3,161 mAh g^-1 at 0.05 C, corresponding to an ICE of 84% and a capacity of 2,550 mAh g^-1 after 100 cycles at 0.2 C. Even at a high rate of 1 C, the capacity is maintained at 1,371 mAh g^-1.

Continuous microporous Si

Different from porous particles made of nano-units, 3D continuous porous Si networks containing continuous nano-skeletons may have better structural stability to alleviate fracture during volume changes. Continuous nanoporous Si microparticles have been demonstrated to improve the cycling characteristics because the connected 3D ligaments disperse the stress during the lithiation/delithiation or sodiation/desodiation of Si. Kim et al. reported 3D porous bulk Si microparticles fabricated by thermal annealing and etching of the composites of butyl-capped Si gels and SiO₂ nanoparticles^[130]. Because the thin pore wall of ~40 nm can accommodate a large strain without pulverization, the porous Si particles maintain a high charging capacity of 2,800 mAh g^-1 at a current density of 2,000 mA g^-1 after 100 cycles, together with excellent rate performance. Jia et al. designed lotus-root-like 3D mesoporous Si by MgR and acid etching^[98]. Porous Si has uniform mesopores and the particles are composed of 30 nm Si crystallites that facilitate fast electrode kinetics. The 3D mesoporous Si has high electrochemical reversibility with a capacity retention of 94.4% at 100 cycles.

Liu and co-workers reported a sub-micrometer mesoporous Si sponge prepared by anodization, and the porous Si showed a capacity of 750 mAh g^-1 based on the total electrode weight with an 80% capacity retention over 1,000 cycles^[102]. In situ TEM demonstrated that the mesoporous Si sponge limits the volume expansion of the active particles and prevents pulverization of the bulk Si particles [Figure 10A]. Choi et al. fabricated porous Si microparticles via an aluminothermic reaction and subsequent MgR [Figure 10B]^[105]. Because of the special porous structure, which can accommodate large volume changes and facilitate ion transportation during cycling, the porous Si/Al₂O₃ microcomposite exhibited outstanding stability with a capacity of 1,500 mAh g^-1 after 100 cycles at C/5 and a small volume expansion of 34% [Figure 10C]. Tao et al. adopted chemical dealloying to prepare a series of porous Si materials by etching Al-Si alloys [Figure 10D]^[131]. The porous Si maintained a spherical structure and the pores were uniformly distributed. As a result, the porous Si had a high capacity of 3,004 mAh g^-1 at 500 mA g^-1 in the first cycle with an ICE of 81% [Figure 10E]. The reversible capacity of porous Si was maintained at 1,525 mAh g^-1 after 100 cycles at 1,000 mA g^-1. This remarkable stability was attributed to the porous structure, which alleviates the volume change during repetitive lithiation/delithiation [Figure 10F].

Figure 10. (A) TEM images and selected area electron diffraction patterns of mesoporous Si sponge particles at different lithiation states. Reproduced with permission^[102] (Copyright 2014, Nature Publishing Group). (B) SEM images of macroporous Si structure. (C) Cycling performance of porous Si/Al₂O₃ composite at C/5. Reproduced with permission^[105] (Copyright 2015, Elsevier B.V.). (D) SEM images of porous Si. (E) Cycling performance of different porous Si samples at 500 mA g^-1. (F) Volume changes can be restricted by porous construction during cycling. Reproduced with permission^[131] (Copyright 2019, Elsevier Inc.).

Mu et al. investigated the impact of commercial Si nanoparticles (C-SiNP) and mesoporous Si microparticles (MP-Si)^[95]. The MP-Si showed outstanding cycling stability and a high reversible capacity of 1,793 mAh g^-1 after 120 cycles at 200 mA g^-1. The Li-ion storage properties benefit from the two-dimensional (2D) nanoplates and abundant mesopores in the Si microparticles, which not only promote immersion of the electrolyte and the transportation of lithium ions, but also buffer the volume expansion and ensure the structural stability of the electrode. Sohn and co-workers prepared 3D porous Si microparticles by alkali and acid etching^[38]. The 3D porous Si electrode showed a high initial charge capacity of 2,919 mAh g^-1 with an ICE of 76.5% and a capacity of 1,222 mAh g^-1 after 200 cycles with an average CE of ~99%. The excellent electrochemical properties of the 3D porous Si suggest that the etching-induced morphological modification, including pore enlargement and reduction of the Si size, is highly effective in improving the cycling characteristics.

Porous ant-nest-like microscale Si (AMPSi) has been produced by a low-cost and scalable top-down approach, involving thermal nitridation of a Mg-Si alloy and removal of the Mg₃N₂ by-product in an acidic solution [Figure 11A and B]^[51]. The high porosity of AMPSi allows the inward expansion of the 3D interconnected Si nano-ligaments during cycling, which gives rise to high structural stability and negligible particle-level outward expansion [Figure 11C]. Therefore, after carbon deposition, the AMPSi and C composites (AMPSi@C) showed a high ICE of 80.3%, reversible capacities of 2,134 and 1,271 mAh g^-1 at 0.1 and 0.5 C (1 C = 4,200 mA g^-1), respectively, and a 90% capacity retention from 20 to 1,000 cycles in a lithium-ion half-cell. The AMPSi@C//NCM full cell exhibited a high reversible capacity of 134 mAh g^-1 at 0.5 C (1 C = 160 mA g^-1) with a capacity retention of 84% for over 400 cycles together with a high energy density of 502 Wh kg^-1. In situ TEM indicated that AMPSi@C remained stable during lithiation and delithiation, with volumetric expansion introduced to the pores to restrict outward expansion [Figure 11D].

Figure 11. (A) Schematic showing preparation of AMPSi and AMPSi@C. (B) Photograph of an ant nest (scale bar = 20 cm). (C) Schematic of lithiation/delithiation process of ant-nest-like microscale porous Si particles. (D) Time-resolved TEM images depicting the lithiation process of AMPSi@C electrode (scale bar = 200 nm). Reproduced with permission^[51] (Copyright 2019, Nature Publishing Group). (E) Cycling performance of Mp-Si with different surface-engineering treatments. Time-lapse TEM images of lithiation of (F) dense Si and (G) Mp-Si@Si@G microparticles. Reproduced with permission^[132] (Copyright 2019, Elsevier B.V.).

Wang et al. designed mesoporous Si microparticles by the thermal disproportionation of low-cost SiO microparticles and subsequent removal of the SiO₂ by-product^[132]. After graphene cage-encapsulated Si skin-sealed mesoporous Si (Mp-Si@Si@G) microparticles were obtained by a surface-engineering strategy by depositing a dense silicon skin onto each mesoporous Si (Mp-Si@Si) microparticle, encapsulation was performed to form a conformal graphene cage. Mp-Si@Si@G showed a high reversible specific capacity of 1,246 mAh g^-1 after 300 cycles at a rate of C/2, which can be ascribed to the optimal specific surface area and suitable porosity of Mp-Si@Si@G, consequently decreasing the electrolyte accessible area, SEI formation and irreversible Li consumption [Figure 11E]. The morphology of Mp-Si@Si@G was monitored by in situ TEM during lithiation and an apparent volume expansion of the dense Si microparticles was observed [Figure 11F]. Cracks appeared after only 32 s and widened with time. At 160 s, the dense Si microparticles fractured abruptly into smaller particles. In contrast, the structures of the Mp-Si@Si@G microparticles were well maintained throughout the whole lithiation process [Figure 11G] and remained intact even after full lithiation because of the sufficient inner pore space and the mechanically strong graphene cage that limited volume expansion.

Si-based alloys have been developed as anode materials for SIBs with a capacity of 954 mAh g^-1[133], and the volume change of NaSi formed from Si is ~244%, which is much smaller than that for Ge-, Sn- and Sb-based anodes in SIBs. In general, the sodiation of Si has different phase behaviors, insertion voltages, and kinetic barriers compared to the reaction with Li. For instance, because of the difference in the ionic radii of Na⁺ (0.97 Å) and Li⁺ (0.68 Å), the insertion of Na into crystalline Si is limited^[134,135]. Xu et al. obtained amorphous Si particles by thermal plasma CVD of silane, which showed reversible Na-ion storage with a capacity of 270 mAh g^-1 at 20 mA g^-1[134]. Wu and co-workers demonstrated the fabrication of amorphous Si-rich interior cores surrounded by a Sn shell (a-Si_SN) via mechanical alloying^[136]. The amorphous sponge-like a-Si_SN has pores [Figure 12A], leading to a reversible capacity of 151 mAh g^-1 after 100 cycles at 36.25 mA g^-1 and a high rate capacity of 143 mAh g^-1 at 725 mA g^-1Figure 12B and C]. Recently, Qiu et al. prepared mesoporous Si microspheres (MSMs) with homogeneously distributed mesopores ranging from 1 to 10 nm via the MgR of SiO₂^[137]. As shown in Figure 12D, the mesoporous Si microspheres with abundant mesopores exhibited a high reversible capacity of 390 mAh g^-1 at 100 mA g^-1 after 100 cycles [Figure 12E] and a high rate capacity of 160 mAh g^-1 at a high current density of 1,000 mA g^-1 [Figure 12F]. The good cycling and excellent rate capability can be attributed to the small particle size and mesoporous structure. However, for the practical application of Si materials in SIBs, several issues still need to be addressed, including structural optimization, electrolyte design, and the development of novel binders^[138-140].

Figure 12. (A) SEM of a-Si_SN. (B) Long-term cycling properties of a-Si_SN electrode. (C) Rate performance of a-Si_SN. Reproduced with permission^[136] (Copyright 2016, Elsevier Ltd.). (D) SEM images of MSMs. (E) Discharge capacity retention of MSMs and Si microspheres. (F) Rate performance of MSMs. Reproduced with permission^[137] (Copyright 2018, Springer Verlag).

Ge-based porous structures for LIBs and SIBs

Ge is in the same chemical group as Si and is regarded as a promising alloying anode material because of its high theoretical capacities [Tables 1 and 2] and suitable operating voltages for LIBs and SIBs. Moreover, the conductivity of Ge is 100 times higher than that of Si. Similarly, Ge anodes also suffer from severe volume expansions, leading to pulverization during lithiation/delithiation. Porous Ge has garnered significant attention since the pores in porous Ge can act as volume expansion buffers to improve the cycling stability. Table 4^{[39,63,141-149]} lists the synthesis strategies and compares the electrochemical properties of different porous Ge anode materials for LIBs.

Yang and co-workers prepared mesoporous Ge microparticles by a mechanochemical reaction of GeO₂ and Mg powders and subsequent etching in HCl^[142]. The mesoporous Ge showed a reversible capacity of 950 mAh g^-1 and a capacity of 789 mAh g^-1 after 20 cycles at a current density of 150 mA g^-1. Jia et al. prepared 3D macroporous Ge (p-Ge) microparticles with a hexagonal-like morphology by MgR and HCl etching^[63]. The p-Ge anode had a high reversible capacity of 1,131 mAh g^-1 at 1 C after 200 cycles and a high rate capability with a capacity of 717 mAh g^-1 at 5 C (1 C = 1,624 mA g^-1). The internal space in 3D p-Ge buffers the volume change of Ge during cycling and preserves the structural stability of the anode materials. Cui and his team synthesized Ge microcubes with a hierarchical porous structure directly on a titanium foil by a simple hydrogen reduction method [Figure 13A]^[148]. The Ge microcubes showed good cycling and rate characteristics due to the voids, such as a high ICE of 91.8% stemming from the high crystallinity of Ge in reversible lithium insertion and extraction, a less adverse side reaction for irreversible Li loss, and unique hierarchical porous structure for easier electrolyte penetration. The Ge microcubes exhibited a high retention capacity of 1,204 mAh g^-1 after 500 cycles at a rate of 1 C, corresponding to a capacity retention of 94.1% from the first cycle, and a good rate of 1,121 mAh g^-1 at a high rate of 50 C [Figure 13B].

Figure 13. (A) SEM images of porous Ge microcubes. (B) Cycling performance of porous Ge microcubes. Reproduced with permission^[148] (Copyright 2015, American Chemical Society). (C) SEM of mesoporous Ge. (D) Cycling performance of Ge (ZnRR) electrodes. Reproduced with permission^[150] (Copyright 2017, Wiley-VCH Verlag GmbbH & Co.). (E) Top view of pGe with 18.05% porosity and one surface Li. In the lower panel, the process of the gradual lithiation is shown with one, four, and twenty Li surface atoms in sequence. Green, white and purple balls represent Ge, H, and Li atoms, respectively. (F) Band gap evolution as a function of the porosity of fully H passivated pGe (dark squares) and pGe with one (red circles) and two (blue triangles) Li atoms. Reproduced with permission^[152] (Copyright 2020, John Wiley & Sons Inc.).

Choi and co-workers synthesized porous Ge by metallothermic reduction using two different metal reductants by the zincothermic reduction reaction (ZnRR) and MgR, respectively^[150]. Figure 13C shows the SEM image of the ZnRR-synthesized Ge particles [Ge (ZnRR)], which have a mesoporous structure with numerous mesopores. With mesopores in the inner parts, the porous Ge provides extra spaces for the expansion of particles and accommodates the large volume change during lithiation/delithiation. The Ge (ZnRR) electrode showed a high ICE of ≈85%, excellent cycling stability at a low temperature of -20 °C, and a high specific capacity of 566 mAh g^-1 at a rate of 0.5 C after 50 cycles [Figure 13D]. The Ge//LiFePO₄ full cell exhibited an ICE value of 98% in the first cycle and retention of 80% at 0.5 C after 50 cycles at -20 °C. Mishra et al. synthesized porous microscale Ge by the hydrogen reduction of GeO₂^[151]. The porous Ge electrode showed a reversible capacity of 1,300 mAh g^-1 after 340 cycles at 1,000 mA g^-1 and a high capacity of 437 mAh g^-1 at 16 A g^-1. The excellent long-term cycling can be attributed to the porous morphology of the Ge electrode, which can maintain the stress induced by Li ions without crushing, provide good electrolyte infiltration and shorten the Li-ion diffusion length. Sosa et al. studied the effects of surface and interstitial Li on the electronic properties of porous Ge (pGe) using DFT calculations and showed that diffusion is easier due to the ion conductivity [Figure 13E and F]^[152]. The porous framework of Ge not only facilitates the fast diffusion of ions and electrons but also buffers the volume change during cycling.

Kohandehghan et al. reported that amorphous Ge can reduce the nucleation barrier of the Na_xGe phase and improve solid-state ion diffusion to enhance Na-ion storage^[153]. Lu et al. also demonstrated that Ge can be readily and reversibly sodiated after turning crystalline Ge nanowires into amorphous ones in the initial lithiation step^[154]. The amorphous Ge nanowires exhibited a 300% volumetric expansion upon sodiation, corresponding to the formation of Na_1.6Ge. Li et al. prepared porous Ge microparticles with a Mn-Ge alloy by selective corrosion^[155]. The porous Ge microparticles showed a broad ligament size distribution in the range between 100 and 300 nm, as shown in Figure 14A. The porous Ge microparticle electrode showed a reversible capacity of 260 mAh g^-1 in the initial cycle and good cycling stability over 50 cycles at C/5 [Figure 14B]. Yi et al. synthesized mesoporous amorphous Ge microparticles by acid etching the Zintl-phase Mg₂Ge at room temperature and studied the Na-ion storage mechanism in meso-porous amorphous Ge (A-Ge) by in situ Raman scattering and ex situ XRD^[156]. As shown in Figure 14C, A-Ge with a rough surface showed a high reversible charge capacity of 600 mAh g^-1, a capacity retention of 90.2% after 50 cycles at 0.2 C, and a rate capacity of 273 mAh g^-1 at 5 C [Figure 14D]. The in situ Raman spectrum in Figure 14E obtained during sodiation/desodiation reveals no distinct peak shift or peak intensity change in C-Ge during the first charging/discharging cycle. This is because Na ions cannot be easily inserted into the interior of C-Ge during discharging. However, the mesoporous A-Ge showed an apparent signal change during charging/discharging. As shown in Figure 14F, the in situ Raman spectrum discloses that the mesoporous amorphous Ge is active with Na in the SIB. The amorphous phase of the Ge anode lowers the sodiation barrier and increases the Na⁺ storage. The mesoporous architecture also provides fast electronic/ionic transportation channels to enhance the rate and provides space for the volume expansion.

Figure 14. (A) SEM images of porous Ge microparticles. (B) Galvanostatic charge-discharge curves of porous Ge microparticles. Reproduced with permission^[155] (Copyright 2019, Elsevier). (C) SEM image of mesoporous A-Ge microparticles. (D) Rate capability from 0.2 to 5 C of A-Ge. In-situ Raman spectrum of (E) crystalline and (F) amorphous Ge during the first discharge/charge curve. Reproduced with permission^[156] (Copyright 2019, Tsinghua University Press).

Sn-based porous structures for LIBs and SIBs

Like Si and Ge, metallic Sn can alloy with Li and Na as an alternative to traditional graphite anodes due to its high theoretical capacity [Tables 1 and 2], high electrical conductivity, moderate operating voltage, and safety. However, the practical application of metallic Sn to LIBs and SIBs is affected by the huge volume change, aggregation of Sn particles during cycling, serious pulverization, and loss of electrical connection with the current collector. The newly formed surface consumes Li by forming the SEI film, and the aggregation of particles leads to poor kinetics, fast capacity decay, and poor cyclability. Direct or indirect methods to alleviate the volume expansion by reducing the particle size and introducing an accommodation space or buffer are currently being studied.

Pore formation can improve the electrochemical properties of electrode materials. However, porous Sn obtained by a simple method still presents significant challenges due to the low melting point and inevitable electrochemical agglomeration of Sn. Song et al. fabricated porous 3D continuous microscale Sn from immiscible Al-Sn alloys^[157]. The microporous Sn anode showed an outstanding initial capacity of 818 mAh g^-1 and a high CE of 99%. The porous structure of Sn still retained the porous framework after first cycling [Figure 15A and B]. Ryu et al. prepared porous Sn consisting of homogeneous Sn particles prepared by vacuum deposition^[158]. The porosity of the Sn exceeded 98% and the porous Sn anode showed a high initial discharge capacity of 1,939 mAh g^-1 and a charge capacity of 714 mAh g^-1 at 0.1 C [Figure 15C]. The reversible capacity of the porous Sn anode changes from 603 to 378 mAh g^-1 at a rate of 1 C after 100 cycles, corresponding to a capacity retention of 63% [Figure 15D]. As shown in Figure 15E, the volume expansion of the porous Sn anode was only 34% during the first lithiation. Cook et al. prepared porous tin (P-Sn) particles by MgR [Figure 15F]^[159]. The P-Sn particles exhibited a high capacity of 578 mAh g^-1 and the maximum capacity remained at 80% after 300 cycles at 250 mA g^-1 [Figure 15G]. The good cycling performance of NP-Sn was attributed to the porous structure that buffers the volume change during alloying. X-ray microscopy showed that the porosity is retained in both the charged and discharged states. As shown in Figure 15H-J, the size distribution of the P-Sn particles was wide (1-20 μm). The pristine particles showed an interconnected network of random mesopores [Figure 15K] arising from dealloying. Figure 15L and M show that the initial random pores and interconnected ligaments still exist even in a fully lithified state with volume expansion (~300%). Therefore, the porous structure not only accommodates Li ions during volume expansion but also retains space for electrolyte diffusion.

Figure 15. SEM images of porous Sn structure (A) before and (B) after lithiation/delithiation process for first cycling. Reproduced with permission^[157] (Copyright 2015, Springer). (C) Charge/discharge profiles of porous Sn anode during the first three cycles. (D) Cycling performance of Sn samples. (E) Cross-sectional SEM images of the as-prepared sample and after 100 cycles of porous Sn anode. Reproduced with permission^[158] (Copyright 2019, John Wiley & Sons Ltd.). (F) SEM images of porous Sn. (G) Cycling performance of porous Sn. (H) Pristine sample. (I) Electrode recovered after one electrochemical cycle at 1.2 V. (j) Fully lithiated sample at 70 mV. (H), (I) and (J) show the magnified subsurface renderings of the highlighted areas in panels (K), (L), and (M), respectively. Reproduced with permission^[159] (Copyright 2016, American Chemical Society).

Metallic Sn with a porous structure is also promising in SIBs. Wang et al. studied the two-step reaction for sodiation of Sn by in situ TEM, which revealed a two-phase reaction mechanism with a modest volumetric expansion of ~60% and a single-phase reaction with a large volume expansion of ~420% after full sodiation^[160]. The large volume expansion/contraction exerted strong mechanical pressure on the Sn nanoparticles, resulting in the production of cracks, serious pulverization of Sn, continuous growth of the SEI, slow kinetics of Na ions, low CE, and rapid electrode failure due to the loss of contact between particles. The porous structure not only provided more space for the volume expansion of Sn but also reduced the ion transfer distance to improve the kinetics. Kim et al. prepared porous Sn (PSn) by the phase-inversion technique for anode materials in SIBs^[161]. The PSn electrode showed a high initial capacity of 1,066 mAh g^-1 and a reversible capacity of 519 mAh g^-1 after 500 cycles at 0.5 C [Figure 16A and B].

Figure 16. (A) Charge-discharge profiles of PSn for 500 cycles. (B) Cycling performance of PSn. Reproduced with permission^[161] (Copyright 2016, Elsevier). (C) Comparison of voltage and specific capacity of a full battery. (D) XRD patterns of porous Sn/Cu anode at different states during the first cycle €. (F) Refined crystal structures of intermediates, NaSn₃, NaSn, Na₉Sn₄, and Na₁₅Sn₄, in discharging process. Reproduced with permission^[162] (Copyright 2018, Wiley-VCH Verlag GmbbH & Co.). (G) In situ XRD patterns of microsized Sn during the first cycle. Reproduced with permission^[164] (Copyright 2016, Wiley-Blackwell).

Wang et al. obtained 3D porous Sn by tuning the redox potential of Cu⁺/Cu with a ligand of thiourea to trigger the replacement reaction between Cu and Sn^2+[162]. The porous Sn electrode showed a high reversible capacity of 700 mAh g^-1 at 2,500 mA g^-1 for 400 cycles, which was significantly better than that of the bare Sn. The full battery of Sn/Cu//Na₃(VO_0.5)₂(PO₄)₃F₂ had high energy and power densities of 311.7 Wh kg^-1 and 182.5 W kg^-1, respectively [Figure 16C], and good cycling performance for 200 cycles without any obvious degradation observed. As shown in Figure 16D and E, the electrochemical reaction was studied by XRD in selected discharging and charging states. During discharging, the new diffraction peak belonging to the cubic phase of NaSn₃ was found in state B (discharging to 0.2 V). After discharging to 0.1 V (state C), tetragonal NaSn with a space group of I41/acd could be observed. Upon discharging to 0.06 V (state D), the characteristic diffraction peaks of Na₉Sn₄ became evident and after full discharging to 0.01 V, the diffraction peaks at 16.5°, 19.0°, 31.8°, 33.3°, 34.7°, and 37.4° corresponded to the final discharge product of cubic Na₁₅Sn₄. As shown in Figure 16(F), metallic Sn was transformed to cubic NaSn₃ and then orthorhombic NaSn in the initial Na⁺ insertion process. Upon further sodiation, the rhombohedral Na₉Sn₄ containing Sn-Sn dimers with a distance of 2.85 Å and five Na atoms oriented in the same direction was formed. Finally, cubic Na₁₅Sn₄ was generated after full sodiation, in which an isolated Sn atom is surrounded by Na atoms in the relatively symmetric 12-fold coordination. In the reverse process, Na₉Sn₄ was gradually desodiated to Sn with four distinctive plateaus of states E-I, as shown in Figure 16D and E. The diffraction peak of Cu₆Sn₅ showed no change during cycling, which is indicative of its nonactivity in SIBs^[163]. However, Cu₆Sn₅ acted as a buffer layer to improve the electron conductivity between the active Sn and Cu foil. It was speculated that the reversible reactions of Sn were Sn↔NaSn₃↔NaSn↔Na₉Sn₄↔Na₁₅Sn₄. Zhang et al. also observed other sodiation phases, such as NaSn, Na₉Sn₄, and Na₁₅Sn₄, by in situ XRD [Figure 16G]^[164].

Sb-based porous structures for LIBs and SIBs

Sb has attracted significant attention as an anode material for LIBs and SIBs because of its high theoretical capacity [Tables 1 and 2], abundant resources, low cost, and environmental friendliness^[165]. Moreover, Sb has high electrical conductivity, reliable ion diffusivity, and low volume expansion (200%) compared to Si. Sb anodes have a safer Li⁺/Na⁺ insertion potential of ~0.5 V, which not only affords stable working voltages in batteries but also shuns the formation of lithium dendrites at low voltages. Sb with a 2D layered structure offers a desirable framework for fast lithium storage. However, the dense Sb still leads to structural fracturing and instability of the SEI layer on the surface of Sb during charging/discharging to accelerate capacity degradation. So far, efforts have been made to overcome these issues. Pore formation in Sb materials can shorten the diffusion length of Li⁺ and Na⁺ and accommodate the volume expansion to improve the electrochemical properties. Typical examples of porous Sb-based materials and their associated electrochemical properties are presented in Table 5^{[161,162,164,166-175]}.

Hou et al. prepared porous hollow Sb spheres by a facile galvanic replacement method^[175]. The hollow Sb spheres showed an initial specific capacity of 629.6 mAh g^-1 and retained a reversible capacity of 627.3 mAh g^-1 after 50 cycles at 100 mA g^-1, corresponding to a capacity retention of 99.6% [Figure 17A]. In addition, such an electrode exhibited a high rate capacity of 435.6 mAh g^-1 at 1600 mA g^-1 [Figure 17B]. The superior electrochemical properties can be attributed to the unique hollow and porous spherical structure [Figure 17C]. The hollow Sb spheres adapted to the volume change and strain during Li⁺ insertion and extraction, and the 3D Sb structure ensured efficient and continuous electron transportation. The large electrode/electrolyte contact area also provided many active centers for charge transfer to enhance the electrochemical kinetics. Rodriguez et al. synthesized 3D Sb chains with a rhombohedral phase by reduction [Figure 17D]^[176]. The 3D Sb electrode exhibited excellent Li storage properties, as exemplified by a highly reversible capacity of 523 mAh g^-1 at 0.5 C after 100 cycles and a retention of 92% [Figure 17E]. Moreover, these materials showed more stable rates and a high specific capacity of 430 mAh g^-1 at 1 C [Figure 17F]. The sodium storage performance of the 3D Sb chains was better than that of Sb/Gr due to the alleviated volume expansion. Meng and co-workers prepared porous Sb with 3D Sb dendrites by a galvanic replacement reaction between Sb³⁺ and Zn powder^[177]. The porous Sb with 3D dendrites showed a high initial charging capacity of 838.6 mAh g^-1 at 0.1 A g^-1, corresponding to an ICE of 79.9% and a capacity retention of 603.2 mAh g^-1 after 124 cycles.

Figure 17. (A) Cycling performance of Sb hollow spheres. (B) Rate performance of Sb hollow spheres. (C) Schematic of lithiation of Sb hollow spheres. Reproduced with permission^[175] (Copyright 2014, American Chemical Society). (D) SEM of porous hollow Sb chains. (E) Cycling performance of porous hollow Sb chains and Sb/Gr. (F) Rate performance of porous hollow Sb chains and Sb/Gr. Reproduced with permission^[176] (Copyright 2019, American Chemical Society).

Hou et al. prepared porous hollow Sb microspheres (PHMSs) for SIBs using a replacement reaction employing Zn microspheres as the template^[174]. The porous shell was metallic Sb and the inner core was a metallic Zn template. Hollow microspheres with porous shells were obtained after removing the Zn templates. The Sb PHMSs had a high initial charge capacity of 634.6 mAh g^-1 at a current density of 100 mA g^-1, corresponding to an ICE of 64.6%. In addition, the Sb PHMSs exhibited a high reversible capacity of 617 mAh g^-1, a retention of 97.2% after 100 cycles at 100 mA g^-1, and excellent rate performance (312.9 mAh g^-1 at 3,200 mA g^-1). The superior electrochemical characteristics stemmed from the unique porous and hollow structure of Sb PHMSs, which could accommodate the strain and facilitate Na⁺ diffusion during repeated alloying/dealloying. Liu et al. studied the effects of the morphology of porous antimony (P-Sb) and the size of Sb particles on the electrochemical properties^[173]. The coral-like P-Sb exhibited excellent electrochemical reversibility after 200 cycles at 100 mA g^-1. The P-Sb electrode could be reversibly cycled at a high current density of 3,300 mA g^-1 and showed a high capacity of 420 mAh g^-1, indicating the fast transfer of sodium ions and electrons. The P-Sb could adapt to the volume change of Na⁺ insertion and improve the capacity retention. The high surface area of the pores allowed for efficient Na⁺/electron exchange to improve the capacity and cycling life of SIBs. The improvement rendered by the coral-like Sb anode arose mainly from the high porosity, and therefore a suitable porous structure plays an essential role in the electrochemical properties of Sb electrodes.

Ma et al. fabricated mesoporous antimony-based Sb₂O₃@Sb with an ultrafine bicontinuous mesoporous structure in the process of dealloying the Mg₉₀Sb₁₀ precursor^[172]. Sb₂O₃@Sb with a mesoporous structure exhibited an ultrahigh specific capacity of 659 mAh g^-1 at 0.2 A g^-1 and excellent cycling stability [Figure 18A]. Operando XRD measurements were executed to study the Na storage mechanism in the mesoporous Sb₂O₃@Sb electrode [Figure 18B]. At the beginning of the first discharge, only crystalline Sb₂O₃ and Sb peaks could be observed. With the progress of the first discharge, the peak intensities of Sb₂O₃ gradually decreased and disappeared, indicating the conversion reaction of Sb₂O₃ with Na to form Na₂O and Sb. In the later stage, the Sb peaks began to weaken and finally vanished. Moreover, the peak of Sb at 28.7° slightly shifted to lower 2θ angles, perhaps due to the insertion of Na into the Sb lattice at the early stage of sodiation. Afterwards, the well-defined crystalline peaks of hexagonal Na₃Sb appeared in the final stage of the first discharge. In contrast, Li et al. reported multidimensional porous Sb nanostructures (Sb NS) synthesized by a facile chemical dealloying strategy^[178]. Such Sb NS anodes exhibited a high reversible capacity of 620 mAh g^-1 and retained 90.2% capacity after 100 cycles at 100 mA g^-1 [Figure 18C]. An ex situ XRD analysis of the Sb-NS electrodes at different charge-discharge states was carried out to further understand the mechanisms of sodium storage [Figure 18D]. In the discharge process, the intensities of the Sb peaks gradually weakened and almost disappeared when the cell was fully discharged to 0.01 V. Moreover, new peaks of 33.5° and 34.3° appeared, corresponding to the formation of the Na₃Sb phase, indicating the conversion reaction of Sb with Na to form Na₃Sb. During the subsequent charge process, the Na₃Sb peaks gradually disappeared, while the intensities of the Sb peaks increased, suggesting the transformation of Na₃Sb into Sb.

Figure 18. (A) Cycling performance of Sb₂O₃@Sb. (B) Operando XRD patterns of mesoporous Sb₂O₃@Sb electrode. Reproduced with permission^[172] (Copyright 2018, Elsevier). (C) Cycling performance of Sb-NS electrodes. (D) Ex situ XRD patterns of Sb-NS electrodes at different charge-discharge states. Reproduced with permission^[178] (Copyright 2019, American Chemical Society).