Phenolic resin-derived hard carbon for sodium-ion batteries: insights and prospects

Abstract

Phenolic resin (PF) has garnered considerable interest as a precursor for anode materials in sodium-ion batteries (SIBs) owing to its high carbonization yield, tunable molecular structure, and well-established synthetic technology. Despite their promise, these materials still face challenges such as low initial Coulombic efficiency, limited rate capability, and inadequate long-term cycling stability. The rational design of high-performance PF-derived carbon anodes necessitates a fundamental understanding of the relationship between their microstructure and sodium storage behavior. In this review, we start from the polymerization and carbonization reaction of PF and discuss the key issues of PF-based hard carbon, along with the sodium storage mechanism. The recent advances in optimizing PF-derived hard carbon are summarized, encompassing the selection of phenolic resin monomers and modification of PF-based hard carbons and their composites. In addition, we offer some perspectives for the design of better PF-based hard carbons for SIBs.

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Zeyu Zhu

Zeyu Zhu received his BS degree in Chemistry from Xinjiang University in 2024. He is currently a PhD candidate at Xinjiang University. His research interest is focused on hard carbon anode materials for sodium-ion batteries.

Jinlin Pan

Jinlin Pan received her BS degree from Guangxi Normal University for Nationalities in 2023. She is currently a PhD candidate in Xinjiang University. Her research focuses on hard carbon anode materials for sodium-ion batteries.

Binghao Wu

Binghao Wu received his BS degree in Applied Chemistry from Taiyuan Institute of Technology in 2023. He is currently a PhD candidate in the College of Chemistry, Xinjiang University. His research interest is focused on hard carbon anode materials for sodium-ion batteries.

Qiang Li

Qiang Li received his PhD degree from Central South University in 2023. He joined Xinjiang University as a Lecturer in 2023 and was promoted to Associate Professor in 2024. His research focuses on electrolyte and electrode materials for advanced sodium-ion batteries.

Weixiang Li

Weixiang Li received his PhD degree from Beijing Normal University in 2021. From 2021 to 2023, he served as a postdoctoral fellow at the Institute of Chemistry, Chinese Academy of Sciences. In 2023, he joined Xinjiang University as a Lecturer and was promoted to Associate Professor in 2024. His research focuses on anode materials and electrolytes of sodium-ion batteries.

Jingui Duan

Jingui Duan obtained his PhD from the Coordination Chemistry Institute of Nanjing University in 2011. He was subsequently awarded a JSPS fellowship at Kyoto University, under the supervision of Professor S. Kitagawa, from 2011 to 2014. In 2015, he joined Nanjing Tech University. He was promoted to a full professor in 2018 and subsequently joined Xinjiang University in 2024. His current research interests are focused on the design and synthesis of porous coordination polymers and membranes.

Ya-Xia Yin

Ya-Xia Yin is a Professor of chemistry at Xinjiang University. She received her PhD from the Beijing University of Chemical Technology in 2012. Prior to Xinjiang University, she worked at the Institute of Chemistry, Chinese Academy of Sciences as an Associate Professor (2012–2018) and a Professor (2018–2023). Her research focuses on nanostructured electrode materials and electrolytes for advanced Li-ion, Na-ion, Li-S, Na-S, and solid-state batteries.

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Introduction

As global efforts to combat climate change and achieve carbon neutrality intensify, advancements in energy storage technologies are rapidly accelerating. The growing demand for renewable energy and large-scale energy storage systems (EESS) has further propelled the development and application of battery technologies. Lithium-ion batteries (LIBs), known for their high energy density and long cycle life, have been widely used in the portable electronics market and are increasingly deployed in EESS applications.^1,2 However, the limited availability of lithium has spurred the search for alternative storage technologies. Sodium-ion batteries (SIBs) have emerged as a promising alternative to LIBs in renewable energy systems, owing to the similar physicochemical properties of sodium to those of lithium and its abundant natural availability.^3,4

In the pursuit of high-performance SIBs, various anode materials have been explored, including transition metal compounds, alloys, non-metallic elements, organic compounds, and carbon-based materials.^5–8 Hard carbon (HC) is regarded as one of the most promising anode materials for SIBs owing to its low working voltage, good reversible capacity, robust structural stability, and cost-effective synthesis.^9–13 Biomass-derived HCs have garnered significant attention due to the abundance of their raw materials, low cost, and environmentally friendly nature. However, their large-scale production is hindered by challenges such as poor batch-to-batch consistency and difficulties in purification. Phenolic resin (PF) is considered as an outstanding precursor for HC due to its tunable molecular structure, mature large-scale production technology, and high carbon yield.^14,15 It is particularly suitable for the industrial production of carbon anodes for SIBs. Nevertheless, PF-derived HCs still exhibit low initial Coulombic efficiency (ICE) and unsatisfactory rate capability, which are mainly caused by the presence of numerous defects and oxygen-containing functional groups. Overcoming these challenges remains a critical research focus for the development of high-performance PF-derived anodes in SIBs. To date, numerous studies have been reported to enhance the sodium storage performance of PF-based HCs through molecular structure optimization, surface coating, heteroatom doping, and the construction of porous structures.^16–18 In this review, we summarize the types and structures of PFs, their pyrolysis mechanisms, and typical synthetic strategies, including monomer selection, pre-treatment, and composite formation for developing PF-derived carbon materials for SIBs. Finally, we provide a perspective on future research directions in the field of PF-based carbon anodes.

In principle, the sodium storage behavior of HC is closely related to its microstructural parameters, including the lateral crystallite size (L_a), stacking height (L_c), interlayer spacing (d₀₀₂), defect intensity ratio (I_D/I_G), and nanopore size. Fast and efficient sodium storage hinges upon the optimization of key structural parameters, which collectively promote sodium adsorption, intercalation, and nanopore filling. Research on the sodium storage mechanism of HC has evolved from the traditional two-stage model to a more refined three-stage model involving adsorption, intercalation, and pore filling (Fig. 1).^19–22 In the first stage (slope region), sodium ions migrate through the preformed solid–electrolyte interphase (SEI) and adsorb onto weakly bonded sites on the carbon surface, displaying fast capacitive behavior. The influence of defects on sodium storage is complex: an appropriate number of defects can provide additional active sites and enhance adsorption, whereas excessive defect density may reduce ICE by promoting electrolyte decomposition and forming a thicker SEI layer. Moreover, high defect concentrations often lead to greater slope capacity at the expense of plateau capacity, while too few defects are unfavorable for sodium ion adsorption. Curvature also plays a significant role in sodium ion adsorption. Moderate curvature facilitates adsorption, whereas excessive curvature reduces adsorption energy and impedes ion intercalation. Additionally, curvature can alleviate the trap effect caused by oxygen defects, enabling reversible sodium adsorption. The second stage (early plateau region) is dominated by faradaic charge transfer reactions. Sodium ions are incorporated into the carbon matrix via diffusion-controlled processes, primarily through intercalation into graphene interlayers. The d₀₀₂ critically affects sodium ion diffusion and storage capacity. During the charge and discharge processes, the d₀₀₂ interlayer spacing undergoes periodic expansion and contraction. Larger interlayer spacing facilitates sodium ion transport, whereas spacing below 0.36 nm hinders intercalation, resulting in slope-dominated behavior. When the interlayer spacing ranges between 0.36 and 0.40 nm, localized structural rearrangements can shift the storage behavior from the slope-dominated to the plateau-dominated. The third stage (late plateau region) is primarily governed by a pore-filling mechanism, wherein sodium ions aggregate into quasi-metallic clusters. This clustering reduces the local positive charge density, thereby weakening electrostatic repulsion between adjacent sodium ions and facilitating the insertion of additional ions into the pore spaces. L_a and L_c significantly influence both the total and plateau capacities. A larger L_a and L_c correspond to a higher degree of graphitization, fewer defects, and, consequently, lower adsorption capacity. This structure also enhances electrical conductivity, leading to superior rate performance. More importantly, optimal values of L_a and L_c facilitate the formation of a closed pore structure, which increases the capacity in the plateau region of the charge/discharge curve. However, excessively large pores may result in incomplete filling by sodium metal or clusters, which can negatively impact overall storage capacity. Therefore, rational structure design of PF can facilitate the development of high-performance precursors. The following sections will elaborate on both pure PF systems and their composite variants.

Fig. 1 Schematic diagram of the sodium storage mechanism of PF-derived HC.

Molecular structure of PF

Polymerization reaction. Since the first synthesis of PF by A. Baeyer in 1872 from phenol and aldehyde under acidic conditions, PF has undergone rapid development and has become an indispensable polymeric material in our daily lives.²³ The synthesis of PF is primarily based on the polycondensation of phenol and formaldehyde catalyzed by either acids or bases and it can be classified into two types according to the catalyst used (Fig. 2): thermosetting phenolic resin (resole) and thermoplastic phenolic resin (novolac).^14,23,24 Resole is produced with a formaldehyde-to-phenol molar ratio (F : P) greater than 1 under alkaline catalysts such as sodium hydroxide or potassium hydroxide. This type of resin contains reactive hydroxymethyl groups (–CH₂OH), which can undergo further cross-linking reactions to form a highly cross-linked and branched three-dimensional network structure. Novolac is synthesized with a formaldehyde-to-phenol molar ratio (F : P) less than 1 using acidic catalysts such as oxalic acid, phosphoric acid, sulfuric acid, or p-toluenesulfonic acid. It exhibits a linear structure with a relatively low molecular weight, and its molecular chains are terminated mainly by hydroxyl (–OH) and phenolic hydroxyl (–PhOH) groups, lacking reactive cross-linking sites. The three-dimensional network of thermosetting resin contributes to the development of porous carbon materials. In summary, due to its tunable cross-linking density and aromatic-rich structure, PF holds significant promise for the preparation of HC materials and demonstrates great potential for application in carbon anodes for SIBs.

Fig. 2 Preparation conditions for thermosetting (a) and thermoplastic (b) PF.

Pyrolysis mechanism of PF. The synthesis of carbon materials from PF requires heat treatment, and a thorough understanding of the chemical reactions occurring during this process is crucial for guiding the modification of PF-based HC. The pyrolysis process is highly complex, involving reactions such as dehydrogenation, condensation, hydrogen transfer, and isomerization, with elementary reactions often occurring simultaneously with other reactions.^24–27 Typically, the reaction mechanisms are inferred by analyzing resin residues treated at different temperatures or volatile species released during pyrolysis. However, this decomposition is not a simple depolymerization process. The pyrolysis of PF generates various small molecules, such as H₂, CH₄, ethylene, methanol, CO, CO₂, and alkylphenols. Some of these molecules can further react with PF, thereby increasing the complexity of the pyrolysis process. Based on previous studies, the pyrolysis of PF can be divided into three stages depending on the temperature (Fig. 3): formation of additional cross-linking, cleavage of cross-links, and polycyclic aromatization. Stage 1 (<350 °C): additional cross-links are formed due to condensation reactions between functional groups such as phenolic hydroxyl groups, or between phenolic hydroxyl and methylene groups. During this stage, H₂O is the main product, resulting in slight weight loss (eqn (1) and eqn (2)). Furthermore, the formation of cyclic ethers such as dibenzopyran (eqn (3)) occurs around 400 °C, and cleavage reactions take place as the temperature increases. This process enhances the cross-linking density of PF, contributing to improved thermal stability and higher char yield. Stage 2 (350–500 °C): gradual cleavage of cross-links occurs. The breakage of these bridging bonds leads to the release of end groups, with phenol and its methyl derivatives being the main volatiles. In addition, cleavage of methylene groups causes damage to the polymer backbone, which is considered the main pathway for PF decomposition (eqn (4)). Stage 3 (>500 °C): as the temperature further increases, additional cleavage of methylene bridges and the removal of hydrogen and oxygen atoms from aromatic structures result in the release of hydrocarbons, carbon oxides, and H₂. However, due to differences in functional groups among PF precursors, significant variations exist in cross-linking density and molecular weight, leading to corresponding changes in pyrolysis mechanisms. Therefore, the analysis must be adjusted according to specific conditions. Based on the above pyrolysis reactions of PF, the basic structural units of alkylphenols and their polycyclic aromatization products determine the initial microstructure of the resulting carbon material. However, PF-based HCs generally possess a large specific surface area and abundant oxygen-containing functional groups, which lead to numerous side reactions with the electrolyte. This often results in the formation of a thick solid–electrolyte interphase, contributing to low ICE and poor long-term cycling stability. Furthermore, the rate performance of PF-based HCs is considerably inferior to that of biomass-based and pitch-based HCs. This is because biomass-derived HCs typically feature a porous structure that facilitates electrolyte infiltration and enhances sodium-ion transport. Numerous strategies have been devoted to enhancing the sodium storage performance of PF-based HCs. A recent review highlights the crucial role of phenolic monomers such as phenol, resorcinol, and phloroglucinol in determining the structure and properties of the resulting carbons.¹⁴ Concurrent studies further demonstrate that the molecular architecture of aldehyde comonomers equally influences the final material characteristics. In this review, we consolidate strategies for improving the electrochemical performance of PF-based HCs through three principal avenues: monomer modulation, precursor modification, and composite formation. Special attention is given to the structural design of PFs, particularly the selection of phenolic and aldehyde types, and its impact on the resulting HC structure and sodium storage capabilities. Additionally, we systematically analyze pretreatment methods for PF and its composites with various materials, providing insight for the rational design of resin precursors.

Fig. 3 The pyrolysis mechanism of PF.

PF-based HCs

Polymer monomers. The molecular architecture of monomers significantly influences the resulting resin structure. Precise monomer design enables tunable control over resin configuration, which constitutes a fundamental advantage of resin-derived HCs as anode materials for SIBs. The electrochemical performance of PF-based HCs is strongly influenced by the crosslinking degree of the synthesized resin. By selecting different types of phenolic compounds, the crosslinking degree can be effectively modulated. Appropriate control of crosslinking facilitates the formation of an ideal HC structure featuring expanded interlayer spacing, enlarged graphitic domains, and reduced defect density, thereby improving the energy storage performance of SIBs. Wang et al. regulated the crosslinking degree of PF by employing various phenolic sources, including phenol, resorcinol, and tannic acid, thereby tailoring the HC microstructure for enhanced sodium storage (Fig. 4a).²⁸ This tunability arises because phenol polymerization involves formaldehyde substitution at the ortho and para positions of phenolic hydroxyl groups, followed by polycondensation. An increase in hydroxyl groups corresponds to a decrease in methoxy substituents (Fig. 4b), which suppresses the hydroxymethylation reaction between formaldehyde and phenols. A decline in sp³-hybridized carbon species associated with methylene bridges signifies the reduced crosslinking. The crosslinking degree decreases in the order: phenol–formaldehyde (P–R) > resorcinol–formaldehyde (R–R) > tannic acid–formaldehyde (T–R). After carbonization, HCs with distinct microstructures were obtained. P–HC displays a highly disordered structure due to pronounced crosslinking constraints, whereas T–HC consists mainly of a graphitic phase with narrow interlayer spacing resulting from weak crosslinking. In contrast, R–R with moderate crosslinking produces predominantly graphitic domains in the resulting R–HC. Research indicates that resorcinol, as a phenolic monomer, confers an intermediate crosslinking degree, enabling precise adjustment of the pseudo-graphitic structure, expanded interlayer spacing, and minimized surface defects. Consequently, the optimized HC delivers a reversible capacity of 334.3 mAh g⁻¹ at 0.02 A g⁻¹, with an ICE of 82.7% (Fig. 4c). It retains 103.7 mAh g⁻¹ at a high current density of 2 A g⁻¹ (Fig. 4d). In comparison, overly crosslinked P–R yields a highly disordered carbon, while weakly crosslinked T–R leads to restricted interlayer spacing and diminished capacity.

Fig. 4 (a) Schematic illustration of diverse resin precursors and corresponding microcrystalline structures of derived HCs. (b) The fitted area ratio of sp³ C and C–O species based on high-resolution C 1s spectra of precursors. (c) The first galvanostatic charge/discharge curves of P–HC, R–HC, and T–HC at 20 mA g⁻¹. (d) The rate performance of P–HC, R–HC, and T–HC. (Reproduced with permission.²⁸ Copyright 2024, Wiley-VCH.) (e) Synthesis process of 3-aminophenol formaldehyde resin (ARF). (f) The discharge–charge curves at the first cycle of different ARF HCs. (Reproduced with permission.²⁹ Copyright 2024, OAE Publishing Inc.)

Furthermore, 3-aminophenol, containing an amino functional group, can be polymerized with formaldehyde (Fig. 4e).²⁹ Reducing the phenol-to-formaldehyde molar ratio increases the crosslinking degree of 3-aminophenol–formaldehyde resin (AFR). Elevated crosslinking promotes carbon materials with a higher carbon yield, wider interlayer spacing and a lower specific surface area. The optimized AFR-HC exhibits a high reversible capacity of 383 mAh g⁻¹ at 0.05 A g⁻¹ (Fig. 4f), excellent rate performance (140 mAh g⁻¹ at 20 A g⁻¹), and an ICE of 82%. Recently, a similar dependence on crosslinking has been observed in acid-catalyzed resins.³⁰ Asfaw et al. synthesized resorcinol–formaldehyde resin via acid-catalyzed condensation polymerization. The optimal sample delivered a capacity of 341 mAh g⁻¹ at 0.01 A g⁻¹ and an ICE of nearly 89% in 1 M NaPF₆ in EC : DEC (1 : 1) electrolyte.

Similar to phenol modification, aldehyde modification is also a common strategy for regulating resin molecules. The steric hindrance effect from the aromatic rings in aldehydes can increase the rigidity of the polymer backbone and the internal free volume, thereby preventing excessive graphitization and promoting the formation of closed pores during carbonization. Xiong et al. successfully prepared a resin precursor with steric hindrance effects by polymerizing resorcinol with cinnamaldehyde (RFSH), which contains a styrenic group (Fig. 5a).³¹ The steric hindrance imparted by the aromatic ring within the RFSH precursor effectively suppresses excessive rearrangement of carbon layers and facilitates the conversion of free volume into closed pores during carbonization. The as-prepared HC anode exhibits a remarkably enhanced discharge capacity of 340.3 mAh g⁻¹ at 0.03 A g⁻¹ and improved rate performance (210.7 mAh g⁻¹ at 1.5 A g⁻¹). Furthermore, increasing the number of aromatic rings in the side groups of aldehyde molecules enhances the crosslinking degree and backbone rigidity of PF, which can promote rigidity in the polymer chains and suppress the rearrangement of carbon layers after carbonization (Fig. 5c). By employing resorcinol–formaldehyde (RF), resorcinol–benzaldehyde (RB), resorcinol–naphthaldehyde (RN), and resorcinol–anthraldehyde (RA) resins, the crosslinking degree decreases sequentially as the number of aromatic rings in the aldehyde increases. Song et al. demonstrated that the benzaldehyde-derived hard carbon (BHC) possesses a highly disordered structure and abundant closed pores, and it delivers a high reversible capacity of 324.7 mAh g⁻¹ at 0.02 A g⁻¹.³²

Fig. 5 (a) Scheme of the synthesis process of RFSH and RF precursors. (b) Schematic illustration of the microstructure of RFSH-1400 and RF-1400. (Reproduced with permission.³¹ Copyright 2024, Wiley-VCH.) (c) Schematic diagram of the internal space of all resins. (Reproduced with permission.³² Copyright 2025, Elsevier.) (d) Schematic diagram of the preparation routes for RG-X samples. (Reproduced with permission.³³ Copyright 2021, Elsevier.)

Given the toxicity of phenol and formaldehyde, which are potentially detrimental to the environment, Ghimbeu et al. recently proposed an innovative approach using plant-derived glyoxylic acid as both a crosslinking agent and a catalyst.³³ This method replaces traditional formaldehyde and eliminates the need for strong acid or base catalysts. The environmentally friendly nature of glyoxylic acid and the catalyst-free synthesis process align with the principles of green chemistry and sustainable development, offering valuable insights for the eco-design of future SIBs and large-scale energy storage applications (Fig. 5d). Beyond glyoxylic acid, the pursuit of sustainable pathways for PF-derived HCs has gained considerable momentum. Researchers are actively exploring the use of biomass-derived phenolic compounds, such as lignin depolymerization products, to replace phenol.²⁸ These efforts not only align with the principles of green chemistry but also potentially lower the overall cost of HC anode materials, facilitating their large-scale application in SIBs. However, the sodium storage capacity of HCs produced by this strategy remains unsatisfactory. Future studies should therefore focus on elucidating how specific modification routes affect the resulting carbon microstructure and electrochemical properties. A summary of resins synthesized from different monomers discussed in this section, together with the structure and sodium storage performance of their derived HCs, is provided in Table 1. The choice of monomers significantly influences the performance of HCs. Phenol serves as the primary phenolic monomer in current research, while investigations have progressively expanded to include resorcinol, phloroglucinol, tannic acid, and aminophenol for tailoring the resin architecture. Comparative analysis of reversible capacity reveals that aminophenol-derived HCs deliver superior performance, attributable to nitrogen-containing functional groups that enhance electrical conductivity and expand interlayer spacing. In addition, resorcinol-resin derived carbons achieve the optimized performance among phenol-, resorcinol-, and tannic acid-resins. This is attributed to resorcinol imparting a moderate cross-linking degree, which allows for precise modulation of the pseudo-graphitic structure with expanded interlayer spacing and suppressed surface defects. However, current characterization of the degree of crosslinking primarily relies on techniques such as Fourier-Transform Infrared Spectroscopy (FTIR), making quantitative analysis challenging. This limitation significantly hinders the establishment of reliable correlations between the crosslinking degree and the HC structure. Note that formaldehyde is the predominant aldehyde monomer, while aromatic aldehydes such as benzaldehyde, naphthaldehyde, and cinnamaldehyde have also been employed to tailor closed-pore structures via steric hindrance effects. Glyoxylic acid, containing both aldehyde and carboxylic acid functional groups, enables catalyst-free resin synthesis. Overall, the current palette of monomeric building blocks remains relatively limited, and the underlying tuning mechanisms are not yet fully established, underscoring the need to explore novel monomers for PF resins.

Table 1 Summary of resins synthesized from different monomers, alongside the structural factors and sodium storage performance of their derived HCs

Entry	Monomers	Solution	d 002 (nm)	L a	L c	I D/IG	S BET (m^2 g^−1)	Electrolyte	Capacity (mAh g^−1)	ICE (%)	Rate performance (mAh g^−1)	Capacity retention (%)	Ref.
1	Phenol and formaldehyde	Alkaline	0.381	∼2.01	∼0.98	1.79	10.79	1 M NaPF6 in diglyme	303.3@0.02 A g^−1	82.1	∼200@500 mA g^−1	—	28
2	Resorcinol and formaldehyde	Alkaline	0.373	∼2.31	∼1.05	1.57	7.34		334.3@0.02 A g^−1	80.3	269.2@500 mA g^−1	82.3 (5000 cycles)@0.5 A g^−1
3	Tannic acid and formaldehyde	Alkaline	0.367	∼2.38	∼1.23	1.33	3.51		240.3@0.02 A g^−1	76.5	∼165@500 mA g^−1	—
4	3-Aminophenol and formaldehyde	Alkaline	0.388	2.64	1.40	0.95	11.2		383@0.05 A g^−1	82	140@20 A g^−1	94 (1000 cycles)@2.0 A g^−1	29
5	Resorcinol and formaldehyde	Acidic	0.370	—	—	1.27	144	1 M NaPF6 in EC : DEC (1 : 1)	281.7@0.02 A g^−1	85	41@1.28 A g^−1	97 (150 cycles)@0.08 A g^−1	30
6	Resorcinol and formaldehyde	Alkaline	0.373	12.48	1.05	1.54	15.11	1 M NaPF6 in diglyme	302.7@0.03 A g^−1	83.7	79.5@1.5 A g^−1	76.2 (1000 cycles)@0.6 A g^−1	31
7	Resorcinol and cinnamaldehyde	Alkaline	0.398	9.02	1.08	2.13	0.85		340.3@0.03 A g^−1	88.5	210.7@1.5 A g^−1	86.4 (1000 cycles)@0.6 A g^−1
8	Resorcinol and formaldehyde	Alkaline	0.377	10.34	1.61	2.51	2.56	1 M NaPF6 in EC : DEC (1 : 1)	311.5@0.03 A g^−1	76.5	—	—	32
9	Resorcinol and benzaldehyde	Alkaline	0.380	9.47	1.5	2.88	1.63		324.7@0.02 A g^−1	75.4	—	—
10	Resorcinol and naphthaldehyde	Alkaline	0.376	10.94	1.64	2.44	2.69		288@0.02 A g^−1	74.2	—	—
11	Resorcinol and anthraldehyde	Alkaline	0.372	12.01	1.70	2.15	3.40		277.9@0.02 A g^−1	77.7	—	—
12	Resorcinol and glyoxylic acid	Acidic	0.399	3.06	1.359	0.69	34.8	1 M NaPF6 in diglyme	325@0.02 A g^−1	88.59	161.4@1.0 A g^−1	93.31 (300 cycles)@1.0 A g^−1	33
13	3-Aminophenol and formaldehyde	Alkaline	0.379	7.45	1.5	3	6.07	1 M NaPF6 in DME	311@0.03 A g^−1	81	<250@2.0 A g^−1	—	34
14	3-Aminophenol and formaldehyde	Alkaline	0.382	11.1	1.56	2.47	5.34		351@0.03 A g^−1	83	263@2.0 A g^−1	91 (2000 cycles)@2.0 A g^−1
15	Phenol and formaldehyde	—	0.398	—	—	1.89	2.42	1 M NaPF6 in EC : DEC (1 : 1) with 5 vol% FEC	309.8@0.02 A g^−1	80.69	∼50@500 mA g^−1	79.6 (140 cycles)@0.02 A g^−1	35
16	Phenol and formaldehyde	—	0.402	—	—	1.96	47.29		300.7@0.02 A g^−1	76.38	∼50@500 mA g^−1	91.9 (140 cycles)@0.02 A g^−1
17	Phenol and formaldehyde	Alkaline	0.389	—	—	2.66	35.3	1 M NaClO4 in EC : DEC (1 : 1)	334.3@0.02 A g^−1	84.7	—	—	36
18	Phenol and formaldehyde	—	0.399	16.22	—	—	3.10	1 M NaPF6 in EC : DEC (1 : 1)	410@0.03 A g^−1	84	—	96 (40 cycles)@0.03 A g^−1	37
29	Resorcinol and formaldehyde	Alkaline	0.386	3.681	0.914	1.57	547.4	1 M NaPF6 in diglyme	367@0.03 A g^−1	88.5	∼110@2.0 A g^−1	92 (300 cycles)@0.03 A g^−1	38
20	Resorcinol and formaldehyde	Alkaline	—	6.20	—	3.10	—	1 M NaClO4 in EC : PC (1 : 1)	310@0.02 A g^−1	84	—	—	15
21	Resorcinol and formaldehyde	Alkaline	0.392	3.82	0.99	—	69	1 M NaPF6 in EC : DMC (1 : 1)	248@0.0186 A g^−1	84	—	—	39
22	Resorcinol and formaldehyde	Alkaline	0.40	10.0	1.34	1.80	52.8	1 M NaPF6 in EC : MEC (1 : 1)	192@0.027 A g^−1	57.5	—	—	40
23	Phenol and formaldehyde	—	0.377	4.93	1.50	1.96	1.82	1 M NaPF6 in EC : DMC (1 : 1)	343@0.03 A g^−1	90.4	73@0.6 A g^−1	90 (150 cycles)@0.03 A g^−1	41

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Modification of PF-based HCs. Modification of PFs provides an effective route for introducing specific functional groups and generating controlled porosity, which is crucial for tailoring the microstructure of resin-derived HCs. While conventional HCs are typically obtained through direct high-temperature carbonization, this process significantly influences critical structural parameters, including interlayer spacing, pore size distribution, and defect concentration. For instance, high-temperature carbonization of PF is an effective strategy for improving the sodium storage performance of HCs. Hasegawa et al. investigated the relationship between structural changes and electrochemical performance of PF-based HC materials carbonized at different high temperatures.¹⁶ The sample carbonized at 1250 °C (PF-HCS-1250) delivered a reversible capacity of 311 mAh g⁻¹ at a current density of 0.02 A g⁻¹ in 1 M NaClO₄ in EC : DEC (1 : 1). As the carbonization temperature increases from 800 to 3000 °C, the interlayer spacing decreases from 0.417 to 0.342 nm. Zhang et al. introduced a 400 °C pre-carbonization step under an Ar atmosphere before high-temperature carbonization at 1300 °C, which altered the connectivity of molecular chains and effectively reduced the oxygen content.³⁴ This approach promoted the growth of L_a from 7.45 to 13.7 nm during subsequent carbonization and accelerated the formation of closed pores, thereby improving the electrochemical performance of resin-based HCs. The HC-400/1300 sample achieved a capacity of 351 mAh g⁻¹ at a current density of 0.03 A g⁻¹ with an ICE of 83%. In comparison, the sample without low-temperature pre-carbonization only exhibited a capacity of 311 mAh g⁻¹ and an ICE of 81%. However, the effectiveness of structural modulation through high-temperature is limited. The reported works with significantly improved performances employed some typical approaches, including pre-oxidation, pore-forming strategies, morphology control, and coating and surface modification. These strategies provide important theoretical foundations and practical guidance for the design of high-performance hard carbon anode materials.

Firstly, pre-oxidation treatment can effectively enhance the cycling stability of HCs by introducing more oxygen-containing functional groups. Yin et al. performed pre-oxidation on the PF precursors,³⁵ obtaining HCs that demonstrated a capacity of 300.7 mAh g⁻¹ at a current density of 20 mA g⁻¹ in an electrolyte composed of 1 M NaPF₆ in EC : DEC (1 : 1) with 5% FEC. Additionally, it exhibited a promoted capacity retention rate of 91.9% after 140 cycles, significantly outperforming the non-pre-oxidized sample (79.6%). This is attributed to the formation of a thinner and more stable SEI film on the pre-oxidized HC surface, facilitating rapid sodium-ion transport and further enhancing cycling performance. Shi et al. investigated the regulation of crosslinking degree in PF precursors by varying the pre-oxidation duration from 1 h to 32 h.³⁶ Pre-oxidation promotes polymerization and crosslinking between molecular chains of PF, forming a highly crosslinked three-dimensional network structure that effectively suppresses the violent decomposition of small molecules during carbonization. The optimized sample exhibited an ICE of 84.7%, representing a 22.2% improvement over the original sample (62.5%), along with a reversible capacity of 334.3 mAh g⁻¹ (at a current density of 20 mA g⁻¹).

Furthermore, pore-forming strategies can regulate the pore structure of HCs, thereby influencing their sodium storage capabilities. Recently, Hu et al. regulated the closed-pore structure of HCs by introducing EtOH as a pore-forming agent in a solvothermal process.³⁷ As the EtOH content increases, the number of closed pores first increases and then decreases, while the specific surface area progressively increases, resulting in a decreased ICE. The optimized material delivered a reversible capacity of approximately 410 mAh g⁻¹ at a current density of 0.03 A g⁻¹, along with a ICE of 84%. Besides EtOH-induced pore formation, the polyvinyl butyral (PVB) templating method offers an alternative strategy. Chen et al. employed PVB as an in situ pore-forming agent to fabricate closed-pore-rich HC from furfural resin.³⁸ The decomposition of PVB produces abundant micropores within the resin. The obtained HC exhibited a high reversible capacity of 367 mAh g⁻¹ and an ICE of 88.5%.

Controlling the morphology can offer advantages, which is illustrated by recent results on PF-based HCs through various strategies. These include modulating solvent polarity in PF synthesis to obtain smaller and more uniformly sized graphitic microcrystals,¹⁵ employing SiO₂ spheres as hard templates to fabricate hollow mesoporous carbon structures (HCSs),^42,43 and incorporating carbon nanotubes into composite architectures. For instance, Ma et al. successfully synthesized HC spheres by carbonizing 3-aminophenol-formaldehyde resin-coated carbon nanotubes (CNTs).⁴⁴ This hierarchical structure, abbreviated as HCS-CNTs, not only establishes the long-range conductive networks that enhance electron transport but also forms a porous framework conducive to ion diffusion. As a result, the HCS-CNT composite demonstrated a capacity retention of 95.1 mAh g⁻¹ after 500 cycles at 1.0 A g⁻¹, highlighting exceptional cycling stability and rate capability. Mai et al. fabricated HC nanofibers (HCNFs) through co-carbonization of nanofibers with PF.⁴⁵ The optimized HCNFs featured the short-range graphitic domains and sufficient interlayer spacing, providing favorable pathways for Na⁺ intercalation. The material delivered a high reversible capacity of 388 mAh g⁻¹ at 0.03 A g⁻¹ and maintained 167 mAh g⁻¹ at 0.5 A g⁻¹, underscoring its outstanding rate performance.

Recently, studies have shown that surface modification and coating techniques (e.g., CVD carbon coating, surfactant-assisted treatment, and graphene oxide-induced graphitization) effectively reduce the specific surface area, improve the ICE, and enhance cycling stability.^39–41 Zhou et al. employed sodium linoleate (SL) as both a surfactant and a catalyst in the synthesis of PF-based HC microspheres (HCMSs) (Fig. 6a).⁴⁶ The addition of SL reduced the average diameter of the PF microspheres (PFMSs), promoted ring-opening polymerization of benzoxazine, and enhanced the crosslinking density. After carbonization, the resulting HC exhibited a smaller particle size (0.236 µm), while maintaining a low specific surface area of 23.83 m² g⁻¹. Job et al. applied a carbon coating layer with graphitic characteristics to resorcinol–formaldehyde resin through chemical vapor deposition.³⁹ This graphitic coating effectively reduced the specific surface area by covering micropores, thereby enhancing the reversible capacity to 248 mAh g⁻¹, alongside an improvement in ICE from 29% to 84%. Huang et al. implemented a graphene oxide (GO)-induced graphitization strategy,⁴¹ where GO templates guided the alignment of phenolic resin molecules during pyrolysis (Fig. 6b). This method effectively reduced defect concentration and specific surface area, yielding a hard carbon material with a high ICE of 90.4% and a reversible capacity of 343 mAh g⁻¹ after 100 cycles at 0.03 A g⁻¹.

Fig. 6 (a) Schematic diagram of the preparation process of PFMS-x with/without SL. (Reproduced with permission.⁴⁶ Copyright 2024, Elsevier); (b) the proposed formation mechanism of GHCs in three stages. Stage I, PF molecules undergo a polymerization reaction (<200 °C). Stage II, some volatile molecules (including H₂O, CO, CO₂, etc.) migrate out along graphene nanosheets (<500 °C). Stage III, the aromatic ring near graphene has a strong tendency to be arranged along the graphene layer during high-temperature carbonization (>1000 °C). (Reproduced with permission.⁴¹ Copyright 2022, Elsevier).

PF-based composites. In addition to monomer modulation and precursor pretreatment optimization, the blending of different precursors can induce cross-linking or facilitate the formation of heteroatom-containing structures. This approach enables the creation of a more diverse closed-pore configuration, ultimately enhancing sodium storage performance. Numerous studies have indicated that introducing small molecules (such as sucrose, triethylenediamine (TEDA), cetyltrimethylammonium bromide (CTAB), 3,4,9,10-perylenetetracarboxylic dianhydride (PTCDA), maleic anhydride, and melamine) to increase the cross-linking density of PF can significantly enhance the electrochemical performance of resin-based HCs.^45,47,48 For instance, Xu et al. co-carbonized sucrose with PF, leveraging the interaction between PF and sucrose to suppress foaming and volatilization of pyrolysis intermediates (Fig. 7a),⁴⁹ thereby reducing the surface area during high-temperature carbonization. The optimized sample exhibited an ICE of 87%, compared to the 74% ICE of PF-derived HCs. Xu et al. utilized PTCDA and PF to tailor the microstructure of HC,⁵⁰ effectively reducing defects and suppressing the growth of graphitic domains during carbonization (Fig. 7b). In this process, PTCDA can be converted into PTCA organic acid, which subsequently undergoes esterification cross-linking reactions with PF. The successful cross-linking between PTCA and PF is corroborated by FTIR spectroscopy. The existence of absorption peaks at 1232.0 cm⁻¹ and 3430.7 cm⁻¹, corresponding to the stretching vibration peaks of C–O–C and C=O of esters, respectively, further confirms the formation of ester linkages. Furthermore, thermogravimetric analysis revealed that the cross-linked product (PPF) exhibited a carbonization yield of 49.4% at 800 °C. This value exceeds the theoretical yield of 40.9%, calculated based on the mass ratio of PTCA and PF, indicating that PTCDA plays a critical role in enhancing the thermal stability of the PF molecule and promoting higher carbonization efficiency. The optimized material exhibited a larger interlayer spacing of 0.394 nm, lower structural ordering, and a smaller specific surface area of 4.8 m² g⁻¹, delivering a reversible capacity of 308.7 mAh g⁻¹ and a high ICE of 77.9%, significantly outperforming the unmodified PF-derived carbon (163.3 mAh g⁻¹, 73.3% ICE). Furthermore, cross-linking with melamine enables nitrogen/oxygen co-doping in HCs, introducing graphitic nitrogen, pyridinic nitrogen, and pyrrolic nitrogen structures. These configurations not only enhance the electrical conductivity of HCs but also create additional active sites, thereby improving overall electrochemical performance.⁵¹

Fig. 7 (a) Proposed formation mechanism of hard carbon materials from sucrose/PF precursors toward an extremely low surface area. (Reproduced with permission.⁴⁹ Copyright 2017, American Chemical Society.) (b) Schematic diagram illustrating the synthesis of a PTCDA modified PF-based carbon heterostructure. (Reproduced with permission.⁵⁰ Copyright 2023, The Royal Society of Chemistry.)

The composite of PF with other polymers also serves as an effective strategy for enhancing the sodium storage performance of HCs.^23,52,53 In particular, cross-linking with pitch allows precise modulation of the HC microstructure, thereby further optimizing its electrochemical properties. For instance, Ma et al. precisely fabricated a HC anode with abundant and suitably sized closed pores (∼0.45 nm) and a nanoscale soft carbon coating through chemical cross-linking between pre-oxidized PF and pitch (10 wt%) (PCHC-10) (Fig. 8a).⁵⁴ The carbon coating effectively reduces open pore defects on the HC, thereby improving its electrical conductivity and ICE. In addition, a series of HC anodes (PCHC-x) with tunable structures were prepared by adjusting the pitch-to-PF ratio. A correlation analysis of their physical and electrochemical properties reveals that closed pore volume and size are the primary factors governing the electrochemical performance of HC electrodes in SIBs. In particular, the PCHC-10 anode delivered a capacity of 359.8 mAh g⁻¹ at a current density of 0.03 A g⁻¹ in 1 M NaClO₄ (EC : DEC = 1 : 1) with 5% FEC additive and exhibited an ICE of 74.8%. This performance significantly surpassed that of the unmodified HC anode (272.6 mAh g⁻¹). Guo et al. further optimized the cross-linking reaction between PF and precursor pitch (PI) by pre-oxidizing PF and PI separately before high-temperature carbonization (Fig. 8b).⁵⁵ Following cross-linking, the L_a value of HC exhibited a slight increase, while the L_c value nearly doubled, leading to the formation of a multi-layered, long-range graphitic microcrystalline structure. Meanwhile, the diameter of the closed pores decreased from 2 nm to 1.6 nm, and their number increased. The optimized HC achieved a high capacity of 401.6 mAh g⁻¹ at 0.02 A g⁻¹ and an outstanding ICE of 89.0% in 1 M NaPF₆ (EC : DEC = 1 : 1) (Fig. 8e), while also delivering the highest reversible capacity of 149.7 mAh g⁻¹ at 1.0 A g⁻¹. Table 2 shows a summary of resin-based composite materials discussed in this section, together with the structure and sodium storage performance of their derived HCs. These results reveal that pitch serves as the predominant additive for blending with PF. This approach effectively reduces defect concentration in the resulting HC and improves its ICE. In comparison, composite strategies incorporating other polymers remain limited and require further systematic investigation. Co-pyrolysis of PFs with small molecules such as PTCDA and melamine introduces heteroatoms and enhances cross-linking density, thereby promoting closed-pore formation. Similarly, the incorporation of inorganic salts has demonstrated their effectiveness in developing closed pores. However, these strategies generally yield resin-based HCs with low ICE (<85%). This limitation highlights the need to develop more modifiers to advance the electrochemical properties of resin-derived HC anodes.

Fig. 8 (a) Schematic diagram of the synthesis route of HC and PCHC-x. (Reproduced with permission.⁵⁴ Copyright 2024, Wiley-VCH.) (b) Scheme of the synthesis process of PF-x and PFPI-13. (Reproduced with permission.⁵⁵ Copyright 2025, Wiley-VCH.)

Table 2 Summary of resin-based composite materials, alongside their structure and electrochemical performance of their derived HCs

Entry	Monomers	Modifier	d 002 (nm)	L a	L c	I D/IG	SSA (m^2 g^−1)	Electrolyte	Capacity (mAh g^−1)	ICE (%)	Rate performance (mAh g^−1)	Capacity retention (%)	Ref.
1	Phloroglucinol and glyoxylic acid	TEDA	—	36.7	12.1	1.63	—	1 M NaPF6 in EC : DMC (1 : 1)	294@0.00744 A g^−1	92	—	—	47
2	3-Aminophenol and formaldehyde	CTAB	0.37–0.40	—	—	2.38	98	1 M NaClO4 in EC : DMC (1 : 1)	388@0.03 A g^−1	53	167@0.5 A g^−1	66 (1000 cycles)@1.0 A g^−1	45
3	Phenol and formaldehyde	—	0.381	—	—	—	236.4	1 M NaClO4 in EC : DEC (1 : 1)	267.1@0.03 A g^−1	46.3	∼100@1.0 A g^−1	—	49
4	Phenol and formaldehyde	Sucrose	0.380	—	—	2.10	2.3		323@0.03 A g^−1	86.4	69.6@1.0 A g^−1	—
5	Resorcinol and formaldehyde	—	0.348	—	—	1.63	100.7		163.3@0.03 A g^−1	73.3	23@0.6 A g^−1	—	50
6	Resorcinol and formaldehyde	PTCDA	0.394	—	—	1.91	4.8		308.7@0.03 A g^−1	77.9	84.7@0.6 A g^−1	90.4 (300 cycles)@0.15 A g^−1
7	Phenol and formaldehyde	—	0.385	—	—	2.04	102.2	1 M NaClO4 in EC : DEC (1 : 1) with 5 vol% FEC	236.3@0.03 A g^−1	31.6	—	—	51
8	Phenol and formaldehyde	Melamine	0.384	—	—	2.27	4.71		319.7@0.03 A g^−1	81.6	180.6@0.6 A g^−1	—
9	Phenol and formaldehyde	Pitch	0.352	4.11	3.20	—	11.4	0.6 M NaPF6 in EC : DEC (1 : 1)	284@0.03 A g^−1	88	90@0.6 A g^−1	94 (100 cycles)@0.03 A g^−1	53
10	o-Aminophenol and formaldehyde	—	0.385	1.57	0.84	1.99	941	1 M NaPF6 in diglyme	286.2@0.05 A g^−1	45.8	104.8@20 A g^−1	—	23
11	o-Aminophenol and formaldehyde	Pitch	0.3743	2.12	1.02	1.88	677		349.9@0.05 A g^−1	60.9	145.1@20 A g^−1	94.5 (2500 cycles)@1.0 A g^−1
12	Phenol and formaldehyde	—	0.349	—	—	1.69	2.32	1 M NaClO4 in EC : DEC (1 : 1) with 5 vol% FEC	272.6@0.03 A g^−1	59.3	51.5@0.6 A g^−1	85.4 (100 cycles)@0.03 A g^−1	54
13	Phenol and formaldehyde	Pitch	0.36	—	—	1.57	1		359.8@0.03 A g^−1	74.8	103.9@0.6 A g^−1	91.4 (100 cycles)@0.03 A g^−1
14	Phenol and formaldehyde	—	0.394	4.00	0.78	1.92	7.3	1 M NaPF6 in EC : DEC (1 : 1)	323.2@0.02 A g^−1	—	<100@1.0 A g^−1	—	55
15	Phenol and formaldehyde	Pitch	0.368	4.81	1.4	2.07	7.9		401.6@0.02 A g^−1	89	149.7@1.0 A g^−1	85.3 (1000 cycles)@1.0 A g^−1
16	2,4-Diaminophenol and formaldehyde	—	0.383	1.60	1.42	1.04	90.75	1 M NaPF6 in diglyme	363@0.05 A g^−1	73	78@20 A g^−1	—	56
17		Zn(C2H3O2)2	0.404	2.12	1.44	0.97	103.68		546@0.05 A g^−1	84	140@50 A g^−1	77 (5000 cycles)@2.0 A g^−1
18	Resorcinol and formaldehyde	—	0.382	2.45	0.86	1.76	23		376.5@0.05 A g^−1	71.7	86@20 A g^−1	71.2 (3000 cycles)@2.0 A g^−1	18
19		ZnCl2	0.406	2.57	1.2	1.72	48.9		501@0.05 A g^−1	77.63	230@20 A g^−1	93.9 (3000 cycles)@2.0 A g^−1
20	3-Aminophenol and formaldehyde	—	0.365	5.71	0.84	1.51	153.9		279.3@0.02 A g^−1	80.3	<100@2 A g^−1	—	57
21	3-Aminophenol and formaldehyde	ZnCl2	0.380	4.58	0.79	1.65	499.8		283.4@0.02 A g^−1	76.5	<100@2 A g^−1	—
22	3-Aminophenol and formaldehyde	Na2C2O4	0.372	6.87	0.81	1.55	68.1		313.3@0.02 A g^−1	74.3	<100@2 A g^−1	—
23	3-Aminophenol and formaldehyde	ZnC2O4	0.375	5.49	0.83	1.57	24.3		349.3@0.02 A g^−1	81.5	221.6@2 A g^−1	95.7 (5000 cycles)@1.0 A g^−1
24	Phenol and formaldehyde	—	0.3356	—	—	0.96	380	1 M NaClO4 in EC : PC (1 : 1)	215@0.1 A g^−1	30	—	—	58
25		Fe(NO3)3·9H2O	0.36–0.40	—	—	1.06	664		245@0.1 A g^−1	33	∼55@5 A g^−1	—
26	Phenol and formaldehyde	PTCDA	0.42	—	—	1.14	56.4	1 M NaPF6 in diglyme	240.8@0.1 A g^−1	82.03	201.3@1.0 A g^−1	100 (100 cycles)@0.1 A g^−1	59
27	Phenol and formaldehyde	PTCDA/Fe(NO3)3/S	—	—	—	1.59	109.5		311.0@0.1 A g^−1	93.89	269.3@1.0 A g^−1	90.8 (100 cycles)@0.1 A g^−1

---

In recent years, strategies including transition metal catalytic graphitization, chemical etching, single-atom doping, and anion/cation synergistic pore creation have enabled precise optimization of the microstructure of HCs.^56–58,60 These approaches effectively expand the graphitic interlayer spacing (to over 0.40 nm), introduce abundant micropores (∼0.8–1.2 nm), reduce defect density, and optimize the SEI. Consequently, the resulting HCs deliver a high reversible capacity, excellent rate capability, a high ICE, and remarkable low-temperature performance. Zhao et al. modified resorcinol–formaldehyde resin through a zinc oxide-assisted bulk etching strategy (Fig. 9a).⁶⁰ Zinc oxide volumetrically etches the graphitic layers via the reaction: ZnO + C → Zn↑ + CO↑. Simultaneously, the resulting zinc can catalyze the local graphitization of the hard carbon. This process resulted in expanded graphitic interlayer spacing and promoted the formation of micropores. The resulting material exhibited the remarkable reversible capacity both at room temperature (RT) and under low-temperature conditions – 501 mAh g⁻¹ at 0.05 A g⁻¹ in 1 mol L⁻¹ NaPF₆ in diglyme electrolyte at RT with an ICE of 77.63% and 426 mAh g⁻¹ at −40 °C.⁵⁷

Fig. 9 (a) Schematic diagram illustrating the synthesis route of modified resorcinol–formaldehyde resin through a zinc oxide-assisted bulk etching strategy. (Reproduced with permission.⁶⁰ Copyright 2025, Wiley-VCH.) (b) Schematic illustration of the formation process of HSC-Fe. (Reproduced with permission.⁵⁹ Copyright 2025, Wiley-VCH.)

Besides zinc salts, iron salts can also be employed to catalyze graphitization, thereby enhancing the sodium storage performance of PF-derived HCs. Xin et al. reported a facile iron-ion catalytic method to synthesize PF-based HC aerogels.⁵⁸ The chelation between metal ions and polymer monomers improves the dispersion of metal catalysts within the carbon precursor, facilitating more effective catalytic conversion of sp³ carbon to sp² carbon. Chen et al. introduced sulfur and iron-based compounds into PF/PTCDA-derived HCs, designating the resulting material as HSC-Fe (Fig. 9b).⁵⁹ This treatment increased the content of oxygen-containing functional groups (C=O) and generated trace amounts of FeS, which promoted reversible Na⁺ storage and accelerated electrochemical reaction kinetics. The optimized HC exhibited a unique layered structure and abundant closed pores. It delivered a high reversible capacity of 311.0 mAh g⁻¹ at 0.1 A g⁻¹ with an exceptional ICE of 93.89% and maintained a capacity of 269.3 mAh g⁻¹ after 800 cycles at a high current density of 1 A g⁻¹.

Summary and prospects

PF has gained significant research attention as a promising precursor for anode materials in SIBs, owing to its high carbonization yield, tunable molecular structure, and well-established industrial synthesis. This review provides a systematic overview of recent progress in PF-derived HC materials, focusing on sodium storage mechanisms, structural modulation strategies, pre-treatment and composites. In summary, PF monomers, the cross-linking degree, and other polymer or small-molecule modifiers have a great influence on the microstructure of resultant PF-based HCs via introducing heteroatoms and defects, modulating the graphitic domain and porosity and consequently affecting Na⁺ storage as illustrated in Fig. 10. In principle, heteroatom doping (e.g., N) can provide more active sites for Na⁺ adsorption, enhance electrical conductivity and expand interlayer spacing, thus increasing the reversible capacity and rate capability. The incorporation of other heteroatoms such as phosphorus and sulfur can be systematically investigated to enrich PF-derived HCs, albeit too much heteroatom doping has the potential risk of more side reactions with the electrolyte. Cross-linking PF with other hydroxyl/carboxyl/aldehyde-containing small molecules or polymer compounds through dehydration, esterification reaction, etc. can promote the cross-linking degree of PF, which is one important factor closely related to the formation of a graphitic domain and nanopores.^{38,47,50,51,53} Bearing in mind that the more graphitic domain is beneficial for promoting the electronic conductivity, the narrow d₀₀₂ spacing (<0.36 nm) of the ordered stacking of graphitic layers makes Na⁺ intercalation thermodynamically unstable.^61,62 The decomposition of the carbonyl/carboxyl group of cross-linked PF could facilitate the formation of porous structures, ranging from macro- and meso- to micro-pores. The micropores are beneficial for Na⁺ accommodation, while macro- and meso-pores expose more surface area to electrolyte permeation and aggravate the electrolyte decomposition reaction. Consequently, a moderate cross-linking degree of PF is deserved with the combination of a suitable pseudo-graphitic structure with expanded interlayer spacing and suppressed surface defects.

Fig. 10 Schematic diagram of the relationship between the PF molecular structure and key characteristics of HC.

PF-derived HC materials show considerable promise as anodes for SIBs. Nevertheless, several key challenges remain to be addressed. Firstly, a primary future direction involves developing an integrated methodology that combines advanced characterization techniques with computational simulations. Advanced in situ multimodal experimental methods include thermogravimetric-infrared-mass spectrometry to reveal the decomposition reaction of PF, solid-state nuclear magnetic resonance to monitor the formation of Na-intercalation compounds or metallic sodium clusters, and cryo-electron microscopy to visualize the SEI formation. On the other hand, theoretical evaluation includes molecular dynamics simulations, and density functional theory calculations can do great favor in clarifying the sodium storage mechanism from the viewpoint of thermodynamics and guide the design of high-performance HCs. Note that most reported capacity concerning PF-derived HCs remains below 350 mAh g⁻¹, less than that of graphite in lithium-ion batteries. The rate capability and long-term stability of PF-derived HCs still need to be improved. Hard carbon with a tunable pore structure, for example, molecular sieve films prepared via a stepwise desolvation method, demonstrates excellent rate capability, maintaining a specific capacity of 224.0 mAh g⁻¹ even at 5 A g⁻¹, along with 82.3% capacity retention at 2 A g⁻¹.⁶³ Likewise, incorporation of metal–organic frameworks or covalent organic frameworks with PFs might be one promising avenue.⁶⁴ Leveraging their well-defined, tunable pore structures and customizable chemical functionalities, these materials can be used to precisely regulate the closed pores, defects, and graphitic domains within the HCs. In addition, the development of eco-friendly PF synthesis and HC production methods represents another significant frontier. Further research is therefore imperative to advance such strategies, including the use of biomass-derived feedstocks, green solvents, and energy-efficient processes, without compromising electrochemical performance. Furthermore, the ICE of resin-based HCs in ester-based electrolytes remains below 90%. Thus, it is essential to develop compatible electrolytes for PF-derived HCs, for example, through solvents or functional additives, to further enhance ICE, rate capability, and long-term cycling stability. Through ongoing innovation in materials and refinements in processing, performance breakthroughs in PF-based HCs are expected, thereby providing substantial support for the development of high-performance and cost-effective SIBs.

Conflicts of interest

There are no conflicts to declare.

Data availability

No primary research results, software or code have been included, and no new data were generated or analyzed as part of this review.

Acknowledgements

This work was financially supported by the Natural Science Foundation of Xinjiang Uygur Autonomous Region (Grant No. 2024D01D03), the Key R&D Program of Xinjiang Uygur Autonomous Region (Grant No. 2024B01009), the National Natural Science Foundation of China (Grant No. 22279147), and the Xinjiang Tianchi Yingcai Project (No. 5105250181p).

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