Hard carbons for sodium-ion batteries and beyond

LIBs, as representatives of rechargeable batteries, have achieved great success and development, being the main portable power source in people's lives over the past decades due to their high energy density. However, with an increasing population and the development of electronic devices and EVs, the demand for LIBs has significantly increased, which will lead to a deficiency of Li resources on the Earth and consequently increasing prices, therefore, LIBs will gradually become unaffordable [18]. Consequently, finding alternative sources to foster new battery technologies is a critical challenge. Due to the similarity of the physical and chemical properties of sodium with lithium (as they are in the same group of elements in the periodic table), SIBs or Na-ion batteries (NIBs) are regarded as one of the most promising candidates to partially replace LIBs.

The working principle of SIBs is same as LIBs, which is called the 'rocking chair' (figure 2): during charge, the cathode desodiates and the Na⁺ ions 'swim' to the anode through the electrolyte and sodiate into the anode, while electrons move from the cathode to the anode along the external circuit. During discharge, the anode desodiates and Na⁺ ions return back to the cathode through the electrolyte, while electrons are transported from the anode to the cathode along the external circuit, and the battery provides power to the external circuit [19].

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The concept of sodium-ion batteries was first established at almost the same time as lithium-ion batteries (1970s for LIBs and 1980s for SIBs) [19–21]. However, SIBs gradually faded out of sight soon after the successful commercialization of LIBs in the 1990s [8], due to their lower specific capacity, lower energy density and worse rate and cycling performance. The main and simplified reason why SIBs perform worse than LIBs is that sodium is larger, heavier, and has a higher standard potential compared to lithium, which causes SIBs to have a decreased capacity and a reduced range of redox potential. The radius and atomic weight of sodium is 1.06 Å and 23 g mol⁻¹, respectively, while lithium is only 0.76 Å and 6.9 g mol⁻¹ [9] (table 1). This makes for larger volumetric changes when the Na concentration changes, which can end up resulting in a strain build-up upon the insertion or extraction of Na⁺ ions and a reduce the cycling life. The larger size of the Na⁺ ions will also affect the kinetic process and cause slower diffusion, which is the main reason for the poorer rate performance [22, 23]. In addition, graphite is usually selected as the anode material in commercial LIBs due to its high cost-effectiveness, and it has a high theoretical specific capacity of 372 mAh g⁻¹ [24] due to the formation of graphite intercalation compounds (GICs). However, the bigger Na⁺ ions do not intercalate into the graphite layers as do Li⁺ ions, due to thermal dynamic issues, which further increases the challenges of finding suitable anode materials for commercial SIBs [25].

Nevertheless, the assumption that Na⁺ ions are too large to intercalate into graphite layers (thereby making graphite a poor material for storing sodium) is an oversimplification of a very complex picture. K⁺ ions with an even larger ion radius (1.38 Å) can happily intercalate into graphite [2, 26]. A graphite lattice can accommodate Li⁺ ions up to a concentration of LiC₆ with a very high theoretical specific capacity of 372 mAh g⁻¹. K⁺ ions can form K-GICs with a theoretical capacity of 279 mAh g⁻¹ for KC₈ [2]. However, in the case of sodium, the concentration is only NaC₁₈₆ [27] or NaC₆₄ [28] with the extremely low theoretical capacities of 12 and 35 mAh g⁻¹, respectively. Rb⁺ and Cs⁺, which have even larger radii, were also investigated for their ease of intercalation into graphite layers [2, 29]. Besides, the desolvation energy for desolving the alkali metal ions from the electrolytes and transferring them to naked ions decreases in most electrolytes from Li⁺ to K⁺ ions, meaning that Li⁺ ionintercalation should 'theoretically' be the most difficult situation compared to Na⁺ and K⁺ ions when forming GICs [30, 31]. This inverted phenomenon results in another controversial opinion that it is not the Na⁺ ions but actually the Li⁺ ions that suffer from the special intercalation situation. The reason for the inferior intercalation behaviour of Na⁺ ions into graphite was calculated to be that the attractive interaction between the Na⁺ ions and the graphite layers is extremely weak, so that there are not enough energetic forces to motivate the intercalation, compared with plating on the graphite surface [32]. Grande et al [33] also found, based on van der Waals density functional theory (DFT) calculations, that the formation enthalpies are positive for NaC₆ or NaC₈ but negative for LiC₆, implying that the formation of Na-rich GICs requires a negative redox potential vs. Na⁺/Na [34, 35].

On the other hand, the intercalation potentials of Li⁺ ions (vs. Li⁺/Li) and Na⁺ ions (vs. Na⁺/Na) into host materials are dependent on the materials' structures. Generally, for common hosts, the standard intercalation potentials are lower for sodium compared to lithium due to the higher potential of sodium than lithium (see table 1) [31]. The same trend was also found in cathode materials, so that the operating voltage of LIB and SIB full cells is dependent on the specific combination of both the positive and negative electrode materials. According to the data reported in the literature [31, 36–40], a number of cathode materials have lower sodiation/desodiation potentials (vs. Na⁺/Na) for sodium than the lithiation/delithiation potentials (vs. Li⁺/Li) for lithium. The reduced values of the positive electrodes are greater than those of the negative electrodes between Li and Na cells. Therefore, the cell voltage of SIBs is usually lower than LIBs, which means the energy density of SIBs is not as high, compared with LIBs, in most cases [31].

Based on the aforementioned shortcomings of SIBs, the development of SIB technology suffered from a long period of silence before researchers started to realise the problem of Li resources and cost. However, when compared to LIBs, the advantages of SIBs can be summarised as follows: (1) there are much more abundant and widely-distributed resources of sodium than lithium. There is only 0.0065% of lithium in the Earth's crust which is mainly distributed in South America, while sodium makes up 2.83% of the Earth's crust and is plentiful all over the world especially in sea water, and therefore much cheaper [7, 41]; (2) aluminium can be used as the current collector for electrodes in SIBs compared to copper in LIBs since aluminium will form an alloy with lithium but not with sodium, which reduces the total cost and weight [31]; (3) the desolvation energy of sodium ions in various organic solvents is around 30% smaller than for lithium ions, which could decrease the charge transfer resistance and increase the electrode kinetics [31, 42, 43]; (4) in addition to Ni, Co and Mn, which were found to be the only electrochemical active transition metals among the layered oxide cathode materials, more choices are available for SIBs such as Cu, Fe, Ti, V, Cr, etc., showing different electrochemical properties and providing more opportunities for the development of SIBs [44]; (5) the use of hard carbon anodes and the differences in the thermal, chemical and electrochemical properties enable more possibilities for electrolyte optimization, to obtain more favourable sodium-based energy storage devices [45, 46]. In addition, similar technology and equipment to that used for LIBs could be used for SIBs, which makes practical and industrial manufacturing easier. These benefits give SIBs great potential for various markets such as low-speed electrical vehicles, home-use energy storage, electronic devices and especially large-scale energy storage, etc., in the future.

Generally, the main key parameters for evaluating the performance of SIBanodes include the specific capacity, energy density (gravimetric and volumetric), rate capability, cycling stability, and the initial Coulombic efficiency (ICE), which strongly depends on the electrode materials, electrolytes, and even the additives.

The theoretical specific capacity Q_t of an electrode material can be determined by Faraday's law [47]:

$$\begin{equation}{Q_t}\left( {{\text{mAh }}{{\text{g}}^{ - 1}}} \right) = \frac{{nF}}{{3.6M}}\end{equation} \tag{ 1 }$$

where n is the number of transferred electrons, F is the Faraday constant (96 485 C mol⁻¹) and M is the molecular mass of the active material. In a practical situation with a galvanostatic charge/discharge status, the specific capacity is calculated by [48]:

$$\begin{equation}Q\left( {{\text{mAh }}{{\text{g}}^{ - 1}}} \right) = \frac{{I\Delta t}}{{3.6{\text{m}}}}\end{equation} \tag{ 2 }$$

where I (mA) is the charge/discharge current, Δt (h) is the charge/discharge duration and m (g) is the mass of the active material. Taking carbonaceous materials as an example, the capacity is affected by the morphology, pore structures, ordering structures, and defects, etc., which are discussed in the next sections.

The gravimetric energy density of a full SIB cell is the integration of the voltage and the specific capacity. The theoretical energy density is calculated based on the following equation [49]: the Gibbs free energy of a standard chemical reaction is the sum of the formulation energy of the reactants and products:

$$\begin{equation}\mathop \sum \nolimits^ {\Delta _r}G = \mathop \sum \nolimits^ {\Delta _f}{G_p} - \mathop \sum \nolimits^ {\Delta _f}{G_r}.\end{equation} \tag{ 3 }$$

The maximum electrical work of the reaction equals Δ_rG:

$$\begin{equation}{\Delta _r}G = - nFE\end{equation} \tag{ 4 }$$

where n is the number of transferred electrons of one mole of reactant and E is the thermodynamic equilibrium voltage. Therefore, the theoretical gravimetric energy density can be calculated as:

$$\begin{equation}{\varepsilon _M} = \frac{{{\Delta _r}G}}{{\mathop \sum \nolimits^ M}}\end{equation} \tag{ 5 }$$

where ∑M is the sum of the mole weight of all the reactants. In practical situations, the energy density is determined as [48]:

$$\begin{equation}\varepsilon = \frac{I}{{3.6{\text{m}}}}\mathop \smallint \nolimits^ U\left( t \right)dt\end{equation} \tag{ 6 }$$

where U represents the voltage window (the volumetric energy density could be calculated in the same way but replacing the mass with the volume). Thus, in addition to increasing the specific capacity of the electrode materials, another way to obtain a high energy density is to elevate the working voltage. That requires a higher voltage of the cathode materials and a lower voltage of the anode materials. The average working voltage of a SIB is the quotient of the integrated area of the charge/discharge profile (the specific energy density) and the total specific capacity. In practice, it usually refers to the voltage at the half of the specific capacity or at the half-integrated area of the galvanostatic charge/discharge profiles. Meanwhile, since a high operating voltage may be beyond the electrochemically stable window of the electrolytes and binders, it is also important to expand the thermodynamically stable potential window of the electrolytes by using proper recipes or adding additives, as well as selecting appropriate binders [50]. Furthermore, it is worth noting that in lab-based experiments, it is usual to calculate the energy density of a full cell based only on the mass of the active materials of both cathode and anode. However, in practical or industrial batteries, the energy density calculation should involve the masses of all the other materials including current collectors, binders, additives, separators and cases, etc. [51]. As a result, controlling the mass of those non-active and packing materials is also very important to improve the energy density.

The rate capability represents the reversible capacity of an electrode material at different current densities. Capacities will typically decrease at higher current rates because the electronic current density is much larger than the ionic current density of the electrolyte and electrodes (and the rate of ion transfer across the electrolyte/electrode interface) [52], meaning that stronger polarisation takes place at high current rates.

The cycling performance relates to the stability of the battery during the long-term charge/discharge process. The lifespan of a battery is defined to be the cycle number at which the capacity retention drops to 80% of its initial capacity [53]. Generally, cycling stability is mainly determined by the interface between the electrode and the electrolyte and by the structural stability of the electrode materials. The electrode/electrolyte interphase is affected by the formation of SEI, which is a passivation layer on the surface of the electrode when the Fermi level of the anode is above the lowest unoccupied molecular orbital (LUMO) of the electrolyte (or the Fermi level of the cathode is lower than the highest occupied molecular orbital (HOMO) of the electrolyte) so that the electrolyte is decomposed to form inorganic insoluble compounds [52, 54]. A stable SEI is beneficial to eliminate further decomposition of the electrolyte during cycling, which ensures the reversible insertion/desertion of Na⁺ ions. The mechanical properties of the SEI are also important in protecting the electrode materials against volume expansion and maintaining the adsorption capacity on the surface [54]. For some anode materials such as alloys, the huge volume expansion during charge/discharge significantly limits the cycling stability, and some strategies including surface coating, and heteroatom doping, etc., would be useful to overcome this phenomenon and improve the cycling performance [19, 55]. Good interface and structural stability is required to ensure a high enough Coulombic efficiency (CE) during the charge/discharge process, and it is found that a high CE of 99.96% is needed for commercial batteries when cycling to over 500 cycles [4].

ICE is the quotient of the initial discharge capacity divided by the initial charge capacity (for the anode half-cell it is the quotient of the initial charge capacity divided by the initial discharge capacity) which is also a significant parameter for SIB technology but gets less attention compared with other parameters. Since the cathode material usually serves as the Na source in a full battery, it should provide Na⁺ ions for reversible cycling as well as the consumption of Na⁺ ions for SEI formation at the first cycle [56]. Therefore, not only does ICE reveal the SEI formation and irreversibility of sodiation/desodiation reactions at the first cycle, but also it affects the energy density of a full battery, as more cathode materials have to be used to cover the initial irreversible Na⁺ ion consumption which results in an increased weight of the battery and a decreased energy density. Some factors were found to influence the ICE including irreversible decomposition of the electrolyte, an irreversible reaction during sodiation/desodiation, and the defects and porosity, etc. As a result, optimising the electrolytes and tuning the nanostructures and surface area would be helpful to improve the ICE [56].

N-doped hard carbons are the most widely investigated heteroatom-doped carbonaceous anode materials for SIBs due to their simple synthesis and fabrication. N-doping in hard carbon materials can increase the electrochemical activity, electronic conductivity, surface wettability and interlayer spacing. N is also able to provide more extrinsic defects meaning that more Na⁺ ions can be adsorbed [122]. Typically, there are three types of doped nitrogen on the carbon lattice, pyridinic-N (N-6), pyrrolic-N (N-5) and graphitic-N or quaternary-N (N-Q), which are shown in figure 7(a). It has been reported that pyridinic-N has a lone pair of electrons which are parallel to the graphene planes, while the long pair of electrons of pyrrolic-N are perpendicular to the graphitic planes, which is beneficial for enlarging the interlayer spacing (figure 7(b)) [123].

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Huang and co-workers [127] used polypyrrole (PPy) as a N and C precursor and carbonised it at 600 °C under N₂ to produce functionalised N-doped carbon nanofibers (FN-CNFs). The weight content of the N was measured to be 13.93% using XPS. The resulting FN-CNFs delivered 172 mAh g⁻¹ at 50 mA g⁻¹ and 87 mAh g⁻¹ at 20 A g⁻¹, with a capacity retention of 88.7% after 200 cycles at 200 mA g⁻¹. The authors claimed that the O- and N-containing functional groups could accelerate the redox reactions on the surface. N-doping could create more topological defects and enlarge the interlayer spacing, which provides more sodium storage sites and a high rate performance. Zhang et al [128] reported the use of bacterial cellulose (BC) together with PPy to synthesise carbon nanofibres within N-doped porous carbon (CNF@NPC) as high-performance anodes for SIBs. Compared with pure CNF, the CNF@NPC shows a much better rate and cycling performance. A PPy and graphene oxide (GO)-derived N-doped carbon nanosheet has also been reported to have the excellent electrochemical properties of 350 mAh g⁻¹ at 50 mA g⁻¹ as an anode in SIBs [129]. Wang and Lou et al [130] have reported a free-standing N-doped nanofiber film, produced by electrospinning and subsequent carbonisation. This material was able to deliver 315 mAh g⁻¹ at 0.5 A g⁻¹ and 154 mAh g⁻¹ at 15 A g⁻¹, indicating a superior rate capability. It also has an excellent cycling stability and the capacity was retained at 210 mAh g⁻¹ after 7000 cycles at 5 A g⁻¹. Guo, Yang and Wang et al [131] reported on a directly carbonized chitin as N-doped amorphous carbon nanofibers (NACFs) for high-performance SIBs. Chitin is the second most abundant biomass in the world and is usually utilised as a nitrogen source for various applications [5, 132, 133]. The NACFs were directly carbonized from 500 °C to 900 °C in Ar. The best sample was carbonized at 700 °C (CC700) showing a high surface area of 369.48 m² g⁻¹ and a high N content of 7.29%. It was found that with an increasing carbonisation temperature, the N atoms were converted into pyridinic and graphitic N. According to the XPS result, the CC700 also contained some oxygen groups such as C = O (quinone-type and O–I) and C–OH (phenol-type, O–II), which also favoured the sodium storage performance. Furthermore, the nanofibrous structure was able to create fast electron transport channels. As a result, the CC700 delivered reversible capacities of 320.6, and 120.6 mAh g⁻¹ at current densities of 50 and 1000 mA g⁻¹. A capacity retention of 85% was obtained at 1000 mA g⁻¹ after 8000 cycles, indicating an excellent rate and cycling performance.

B-doping is another substitutional doping strategy for enhancing the electrochemical properties of hard carbon anode materials in SIBs by introducing a p-type doping character to reduce the Fermi level [120, 134]. Cao and co-workers [135] found that every boron atom can adsorb three Na atoms due to the hybridized orbitals with carbons. The sodium adsorption energy on B-doped graphene was calculated to be 2.7 and 7.1 times of the intrinsic energy of N-doped graphene, respectively, via the first-principles density functional theory (DFT), indicating much more stable sodium storage on B-doped carbons. Barone et al [136] have also studied the intercalation of Na⁺ ions into BC₃ (Na_xBC₃) by dispersion-corrected DFT, showing that the most stable intercalation compound for Na is NaBC₃ which corresponds to a theoretical capacity of 572 mAh g⁻¹. The sodiation energy of B-doped graphene is calculated to be ∼0.92 V. Boron is a p-type dopant which prefers to accept electrons from the Na⁺ ion, which means it confers a stronger and stable binding between the doped B and the Na⁺ ions [80]. However, although the B-doped carbons exhibit great potential as sodium storage electrode materials, there are still not many reports about B-doped hard carbons for SIBs since the fabrication and synthesis methods are relatively harder, compared to other heteroatoms. It was also found that B-doping will introduce high-energy in-plane defects and cause significant irreversibility during the first desodiation process [80], which requires specific consideration in the future.

S- and P-doping have also been proved to have positive effects on sodium storage for hard carbon anode materials. Unlike N and B, S and P have much larger covalent radii and can fit between the graphene layers in most cases, which will significantly increase the interlayer spacing. In addition, the S and P can not only enhance the electronic conductivity, but also reversibly adsorb Na⁺ ions and provide additional sodium storage capacity [120, 121].

The most commonly used sulphur source is elemental S, which can be easily mixed and carbonized together with hard carbon precursors. An experimental and theoretical study of S-doped carbon nanofibers (S-CNFs) made by simply mixing bacterial cellulose (BC) with elemental S with a mass ratio of 1:2 and carbonising at 500 °C in Ar was reported by Jiang and Wang et al [124]. The S-CNFs showed a high S content of 17.44 wt.% and an excellent rate capability of 355 and 257 mAh g⁻¹ at 0.05 and 8 A g⁻¹, respectively, with reversible electrochemical reactions between the –C–S–C– bonds with the Na⁺ ions. Theoretical calculations indicated that the interlayer spacing was enlarged and the diffusion barrier was reduced due to the S doping (figure 7(c)). In addition, the deformation charge density (DCD) and density of states (DOS) in figures 7(d)–(e) show that S-doping also improves the electronegativity and electrochemical activity as well as providing more defects, which results in a smaller band gap, which are all beneficial for enhanced sodium storage performance. Biomass-derived S-doped carbon nanosheets (S-CNs) were also synthesised by direct carbonization in Ar upon mixing a biomass precursor (spring onion peel) and elemental S. The resulting S-CNs can deliver ultrahigh reversible capacities of ∼601.2, 545.2, 220.1 and 133.6 mAh g⁻¹ at current densities of 0.05, 0.1, 5 and 10 A g⁻¹, respectively, which are much higher than those of undoped samples. A high capacity of ∼211 mAh g⁻¹ at 5 A g⁻¹ after 2000 cycles was obtained, indicating a great cycling performance [137].

It has been reported that elemental S would lead to unstable configurations in the doped carbon layers, which would make the electrochemical reactions similar to those of Na-S batteries [138]. Other than elemental S, some other S sources have also been investigated for use as a dopant to produce S-doped hard carbons for SIBs. Huang and co-workers [139] pyrolyzed prepared poly(3,4-ethylenedioxythiophene) (PEDOT) at 700 °C in Ar and obtained S-doped carbon (SC) with an enlarged interlayer spacing and a high content of S (15.17 wt.%). Owing to its small surface area of ∼40 m² g⁻¹, the ICE (73.6%) was relatively higher than for other heteroatom-doped carbons. As an anode in SIBs, SC shows excellent stable reversible capacities of 327.8 and 119.5 mAh g⁻¹ at 0.5 and 5 A g⁻¹ with a retained capacity of 322 mAh g⁻¹ at 0.5 A g⁻¹ after 700 cycles. Hou and co-workers [140] used 2-thiophenemethanol as a S source to synthesise with durian shell-based activated carbon. The resulting S-doped carbon can have a remarkable capacity of 100.02 mAh g⁻¹ at 5 A g⁻¹ after 4500 cycles. CS₂, Dimethyl disulphide, Na₂S, H₂S, and SO₂ can also be the S sources for the fabrication of S-doped carbons by either being mixed with the carbon precursors or calcinating the precursors in a S-containing gas atmosphere [141]. S-containing precursors have been found unable to provide sufficient doping, while the major advantage is to increase the d-spacing [138]. S-containing gases cause environmental issues and the doping level is hard to control. In addition, the doped S in the carbon materials is usually located at the defects and edges in a thiophene-type structure [120], therefore selecting proper S sources and synthesising the S-doped hard carbons with a homogenous S distribution would be important aims.

Phosphorus, with low electronegativity and high electron-donating properties can provide more electrochemical active sites and extrinsic defects, and has also attracted enough attention as an alternative doped heteroatom to enhance the sodium storage performance of hard carbons. A comparative study of P-, S- and B-doped hard carbons (P-HC, S-HC and B-HC) was made by using H₃PO₄, H₂SO₄ and H₃BO₄, respectively, with the same doping level of 5 wt.%. Consequently, the P-HC showed more defects and curvatures that may be present because the doped PO_x is not coplanar with the graphene layers, and it delivered the highest specific capacity of 359 mAh g⁻¹ at 20 mA g⁻¹ compared with other HCs (figure 7(f)) [80]. Ji's group [142] has reported on P-doped large-area carbon nanosheets (P-CNSs) obtained by the conversion of 0D carbon dots (CDs) to 2D nanosheets doped with P. The P-CNSs obtained showed a P content of 1.39% and delivered a high reversible capacity of 328 mAh g⁻¹ at 0.1 A g⁻¹. This material also had an excellent cycling stability of 149 mAh g⁻¹ at 5 A g⁻¹ after 5000 cycles. Recently, Liu, Yang and Lu et al [143] synthesised a 3D well-ordered porous P-doped carbon with phosphoric acid as the P source, which exhibits excellent specific capacities of 270 mAh g⁻¹ at a low current density of 0.2 A g⁻¹ and 140 mAh g⁻¹ at a very high current density of 10 A g⁻¹. The extraordinary sodium storage behaviour was concluded to be due to the larger distance between the graphitic layers, the modified electronic structures and the enhanced Na⁺ ion adsorption and diffusion kinetics.

On the other hand, it was also reported that phosphorus is an n-type dopant that is reluctant to accept the electrons from Na, which leads to an unstable binding energy between the doped P and the Na [80]. The large radius of phosphorus usually makes the doping more difficult compared to other heteroatoms [120]. Therefore, consideration of both the advantages and disadvantages would be necessary to design proper P-doped hard carbons in the future for high-performance SIBs.

Doping with oxygen is usually not regarded as a heteroatom-doping strategy for hard carbon materials because there is always a great amount of intrinsic O contained in the hard carbon precursors. These O-containing functional groups will inherently be retained after carbonisation and in most cases, it is not necessary to dope additional O through other routes. These O-containing groups in hard carbon materials have a strong impact on sodium storage behaviours, which cannot be ignored and require more attention to understand the sodium storage mechanisms and design high-performance hard carbon anodes.

The capacity contributed by the voltage range over 0.1 V (vs. Na⁺/Na) is believed to correspond to the reactions between the Na and O in most O-doped carbons [144]. Song and Jia et al [125] prepared N-doped carbon nanospheres with different O contents and observed the effect of the O-containing groups on the sodium storage. As shown in figure 7(g), the authors found that the C = O groups are very stable and the sodiation/desodiation is mostly reversible between the Na⁺ ions and the C = O groups. The C–OH, C–O–C and COOH groups were found to be able to provide more defects and enlarge the interlayer spacing. A higher content of such groups led to higher reaction activity and a higher Na adsorption capacity. Nevertheless, these types of O-containing functional groups also caused strong irreversibility. In a recent study proposed by Huang and co-workers [126], the C = O bond was reported to improve the adsorption capacity of hard carbons in the case of buckwheat husk-derived hard carbons (BPC). BPC carbonised at 1100 °C showed the best performance of 400 mAh g⁻¹ and 116 mAh g⁻¹ at current rates of 0.05 and 2 A g⁻¹, respectively, with a capacity retention of 96% at 2 A g⁻¹ even after 3000 cycles. It was found that with an increase in the carbonisation temperature from 700 °C to 1300 °C, the ICEs of the resulting BPCs also increased from 42% to 72%. Correlated with the ratios of C–O and C = O groups in the BPCs carbonised at different temperatures shown in figure 7(h), the C = O content increased while the C–O increased, suggesting that the C–O groups will cause irreversible loss of capacitiy in the first few cycles.

As opposed to the monoheteroatom doping of carbons, multiatom doping is able to combine the advantages of the single heteroatom with increased electrochemical activity in hard carbon anode materials in SIBs. Jiang and Wang et al [145] synthesised 2D N/B dual-doped carbon nanosheets (NBTs) from gelatin and H₃BO₃. NBTs show pseudocapacitive-dominated sodium storage properties and the performance was much better, compared with samples containing N doping only. The same group also reported on a N/S co-doped hard carbon using sodium citrate and thiourea. As a comparison, undoped carbon from sodium citrate only and a N-doped carbon from sodium citrate and urea were also produced. It was found that the dual-doping significantly enhanced the electrochemical properties of the hard carbon compared with undoped and single N-doped carbons [146]. Zhou and co-workers [138] have reported on a N/S dual-doped carbon nanosheet for SIBs and studied its performance with and without sulphur doping. With the S doping, the material had a larger interlayer spacing and more defects, presenting a better rate and cycling performance compared to the carbon nanosheet that was only N doped. The sodium storage performance of N/P co-doped carbon microspheres (NPCM) was studied by Li and Peng et al [147]. With heteroatom-doping, the NPCM showed a decreased surface area, more defects and enlarged interlayer spacing. In addition, owing to the spherical morphology, the NPCM delivered a much higher specific capacity and rate capability, compared with undoped carbon.

In general, heteroatom-doping is indeed a popular strategy for enhancing the electrochemical properties of hard carbons by increasing the electronic conductivity and wettability, changing the Fermi level, enlarging the interlayer spacing, and providing more extrinsic defects and functional groups for more sodium adsorption sites. Heteroatom-doped hard carbons usually have highly surface-controlled and capacitive capacities and exhibit only sloping regions in their voltage profiles, ensuring higher safety, fast rate capability and stable cycling performance. However, on the other hand, the introduced defects, heteroatom functional groups and the large surface areas from the low-temperature calcination also cause serious irreversibility during the first cycles and thereby a low ICE. In addition, the synthesis methods of the heteroatom-doped hard carbons are relatively more complicated compared with undoped ones, which may increase the cost and limit the commercialisation process of this kind of material for next-generation SIBs. On the other hand, in some cases, it is still not very clear how each heteroatom specifically contributes to the sodium storage performance and there have been some conflicting reports. A deeper understanding of the exact role of each heteroatom must still be established in further works.

A sodium-ion capacitor (SIC) is an energy storage device consisting of a battery-type anode and a capacitor-type cathode, leading to a balance between high-energy sodium-ion batteries and high-power supercapacitors. This occurs because the diffusion-controlled Faradic reactions in the body phase of the anode and the surface-controlled non-Faradic reaction on the surface of the cathode kinetically constrain each other [159]. In most cases, SICs employ metal oxides as anode materials and activated carbons as cathode materials, which are both usually fabricated using critical materials such as Co, Ni and fossil fuels. Moving towards a more eco-friendly and low-cost energy storage system, hard carbons from biomass have been widely investigated as electrode materials for sodium-ion capacitors [160]. Compared with metal oxides, low-cost hard carbons can also exhibit excellent capacities when they are employed as anodes because of the special storage features of sodium ions which have been mentioned above. When it comes to the cathode, although activated carbons from fossil fuels have been considered to be the default choice, porous hard carbons activated from biomass, in fact, possess more advantages including low-cost, tuneable porosity and surface chemistry. The state-of-the-art rational design strategy of hard carbons for both anode and cathode materials has shown promising application in sodium-ion capacitors.

Titirici and co-workers [48] utilized cellulose as a precursor to rationally synthesize both anode and cathode materials, thus assembling an all-cellulose-based sodium-ion capacitor (figure 10(a)). The anode materials were prepared by the carbonization of cellulose nanocrystals and then formed a non-porous but disordered structure with a reversible sodium storage performance of 256.9 mAh g⁻¹ at 0.1 A g⁻¹, while for the cathode part, carbon materials with hierarchically porous structures (2137 m² g⁻¹) which were activated from cellulose microfibrils using KHCO₃ and Melamine co-activation were able to exhibit a discharge capacitance of 212.4 F g⁻¹ at 0.1 A g⁻¹. This project reasonably took advantage of the inherent properties of the hierarchal building blocks of cellulose to achieve a sustainable sodium-ion capacitor with an optimized electrochemical performance. Utilizing Na⁺ and ClO₄− as charge carriers, a full Na-ion capacitor with two asymmetric carbon electrodes was able to reach an energy density of 181 Wh kg⁻¹ at a power density of 250 W kg⁻¹ (figure 10(b)). In addition, Ding et al [157] prepared carbon materials with specific properties from different parts of peanut shells by corresponding methods, based on their understanding of the natural characteristics of peanut shells. For instance, the inner shells of peanuts, made of lignin with disordered repeating units and three-dimensional structures, can lead to an amorphous carbon whose distance between graphene layers is expanded after carbonization. This specific structure can reduce the reaction barrier to Na ion insertion and achieve an enhanced electrochemical performance of 315 mAh g⁻¹ at 0.1 A g⁻¹. However, the outer shell of the peanut is mainly composed of cellulose microfibrils which can be converted into a graphene-like analogue through hydrothermal treatment and chemical activation. The porous carbons achieved, with a high surface area of 2396 m² g⁻¹, delivered a specific capacity of 161 mAh g⁻¹ at 0.1 A g⁻¹. The assembled sodium-ion capacitors were able to deliver an energy density of 201 Wh kg⁻¹ when the power density was 285 W kg⁻¹ (figures 10(c)–(d)). Therefore, sodium-ion capacitors become the trade-off between batteries and supercapacitors.

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Although recent research has demonstrated some successful examples of high-performance sodium-ion capacitors, the drawback of fabricating the full cell configuration still needs further consideration because the pre-sodiation of anode materials is usually carried out in half cells against metallic sodium. Without a pre-sodiation process, the lack of charge carriers would cause the loss of reversible capacity and suppress the cycling performance. The most important research direction, therefore, is the in situ pre-sodiation of anode materials during the material preparation step [161]. To sum up, the successful application of hard carbons has explored the road to future commercialization of sodium-ion capacitors, but more scale-up experiments in pouch cells and deep study of their working principles are needed.

Dual-ion batteries (DIBs) are regarded as one of the most promising alternatives to current Li-ion cells due to their comparable energy density. The working principle of DIBs is different from LIBs/SIBs, though. Hu et al [158] reported an all-carbon DIB working by a simple principle of having an electrolyte sandwiched between a hard carbon anode and a graphite cathode. DIBs show the feasibility of designing an all-carbon energy storage system delivering a specific capacity of over 60 mAh g⁻¹ and a good cycling performance. Upon charging, Na⁺ cations migrate to the anode and are inserted into the hard carbon, while PF₆ ⁻ anions intercalate into the graphite; upon discharging, Na⁺ and PF₆ ⁻ ions move back to the electrolyte (figure 10(e)). The wide range of hard carbon families whose plateau region of both charge and discharge are usually below 0.1 V are suitable for building up a DIB of high energy density, thus hard carbons are categorised as a promising cation hosting material for DIBs. An all-carbon configuration can be achieved in a cell paired with a graphite cathode for anion intercalation (e.g. Hard carbon |1.0 M NaPF₆ in a carbonate electrolyte | graphite). Such a carbon sandwich eliminates the use of any layered-metal oxides. The excellent thermal resistance of the carbon electrodes ensures the safety of all-carbon DIBs. The cathode materials in SIBs, typically exhibit a working voltage of ∼3.5 V, while DIBs can deliver an average working voltage of around 4.5 V. The increase in the working voltage may compensate for the limited capacity graphite cathode which has a theoretical capacity of 112 mAh g⁻¹, resulting in a graphite intercalation compound of C₂₀PF₆ [162]. This strategy of assembling DIBs has been attracting increasing attention in the battery community in recent years and may also extend to multi-ion batteries in the near future with the hard carbons playing an important role [163].

Anode and cathode materials in conventional SIBs can benefit from hard carbon. The NASICON Na₃V₂(PO₄)₃ is among the most promising cathode materials for SIBs due to its conductive framework, large channels for transferring Na⁺ ions and excellent thermal resistance [167, 168]. The rate performance of Na₃V₂(PO₄)₃ is not ideal though, so coating hard carbon onto a Na₃V₂(PO₄)₃ core via a hydrothermally assisted sol-gel approach is one of the solutions to overcoming this disadvantage, improving the electrochemical performance [169–171]. Zhao and co-workers [170] applied conductive polyaniline (PANI) as a hard carbon coating via a facile sol-gel process onto the NASICON Na₃V₂(PO₄)₃ to form a composite. The hard carbon layer consisting of localised graphitised layers was an excellent conductive agent helping charge transfer [61]. After modification, the surface exhibited higher conductivity and protected the Na₃V₂(PO₄)₃ core. The double carbon-coating layer enhanced the conductivity and prevented the Na₃V₂(PO₄)₃ core from growing over cycling. The protective layer also helped to suppress the volume change during charging and discharging. The modified cathode delivered a discharge capacity of 111.6 mAh g⁻¹ at 1 C, 61.0 mAh g⁻¹ at 50 C and good long cycling performance over 15 000 cycles [170].

Regarding anodes in SIBs, hard carbons coupled with transition metal chalcogenide anodes is a trendy topic. Transition metal chalcogenides (TMCs), usually be denoted as MX₂ (M = Ti, V, Nb, Mo, W, Re; X = S, Se, Te), are emerging energy storage technologies, since a weak van der Waals (vdW) force among their interlayers, owing to the layered structure of TMCs, is beneficial for Li⁺, Na⁺, and K⁺ ions to de-/intercalate [172]. TMCs usually have a higher theoretical capacity, for example, MoS₂ is a promising anode for SIBs not only because of its high theoretical Na⁺ storage capability but also its ample space between 2D layers for accommodating more abundant ions [173]. However, TMC materials are not yet commercially possible, as their working principle (conversion reactions), volume expansion and unstable electrochemistry primarily hinder their electrochemical performances for charging and discharging. Therefore, the cycling and rate performance of TMC anodes cannot satisfy the current demands of the market. Hard carbons can be applied to restrict the volume expansion of TMCs, and hydrothermal carbonisation in this regard is an effective way to synthesise hard carbons. Ji and Zhang et al [174] constrained exfoliated MoSe₂ inside carbon spheres with a diameter of 100 nm, building a hybrid structure. Polyvinyl pyrrolidone (PVP) was used as a carbon precursor which can easily be dissolved and chemically attached to the surface of MoSe₂. Due to such a unique carbon layer, C–O–Mo bonds anchored on the surface, which was reported to activate the Se element generated during cycling and lower adverse reactions, promoting capacity, stability and rate properties. Such beneficial bonds are further confirmed by DFT models. Thanks to the carbon layer, the electrochemical results for modified MoSe₂ in SIBs are better than for bulk MoSe₂; after 120 cycles, the specific capacity retained was around 441 mAh g⁻¹ at 1 A g⁻¹ and 348 mAh g⁻¹ at 4 A g⁻¹.

Besides acting directly as active materials, hard carbons can also serve as a functional host for other kinds of active materials with different storage mechanisms including alloying and conversion. For example, hard carbons with stable chemical and physical properties can restrict the volume expansion of metal anodes by allowing the homogenous plating and stripping of metallic sodium on the surfaces of hard carbons [175]. Metallic sodium has been considered to be the ultimate anode material for sodium ion storage because of its high theoretical capacity of 1166 mAh g⁻¹. However, during the charge and discharge process, sodium dendrites will be induced because the metallic sodium plating/stripping is a random and uncontrollable process, thus leading to a quick fading of batteries. Luo et al [164] prepared 3D hard carbon hosts with porous channels derived from wood to encapsulate the metallic Na anodes as well as suppress their dendrite formation (figure 11(a)). Pore and defect structures can induce the nucleation of metallic sodium and help it to plate and strip homogenously. Employing Na-hard carbon composite anodes, the cycling performance of symmetric half cells can be greatly improved, up to 500 cycles at 1.0 mA cm⁻² (figure 11(b)). Zhu et al [176] also prepared 3D scaffolds from lignin with boron doping for metallic anode protection. Heteroatom-doping can form extrinsic defects on the graphene layers which can act as active sites for the initial nucleation of metallic anodes, thus further improving the homogenous plating and stripping of metallic anodes.

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Phosphorus (P) is another promising anode with an ultra-high theoretical capacity of 2596 mAh g⁻¹, but its intrinsically low conductivity and huge volume expansion in electrochemical reactions severely hinder its application in sodium battery anodes [177]. Hence, hard carbons with good conductivity can be utilized as a functional host for phosphorous. Oh and Lee et al [165] prepared P/hard carbon composite anodes by ball milling methods (figures 11(c)–(d)). A reversible capacity of 1890 mAh g⁻¹ was obtained at 0.286 A g⁻¹ since hard carbons can not only create conductive channels for composite anodes but also constrain the excess volume expansion of P (figure 11(e)). The composite anode also exhibited an excellent rate performance at a 10-fold current density, where the reversible capacity was able to remain at 1540 mAh g⁻¹. Without hard carbons, bare P anodes only delivered a reversible capacity of around 400 mAh g⁻¹ with a low ICE of below 25%. Furthermore, in the second cycle, the reversible capacity faded to only half of the previous cycle (figure 11(f)). This poor performance further proves that hard carbons play an important role in the improvement of the electrochemical performance of P anodes.

Hard carbons can also contribute to various advanced cathode materials for next-generation sodium batteries. One of the next-generation sodium batteries is sodium-sulphur batteries where S is used as a cathode material, while metallic sodium is used as the anode material as well as the sodium source in full cells. The reactions between the S and Na⁺ ions can be divided into several steps including the formation of Na₂S₈, Na₂S₄, Na₂S₂ and Na₂S. The highest theoretical capacity of S cathodes can be 1675 mAh g⁻¹ when Na₂S₂ and Na₂S form. However, insulated S has a low reaction activity without any conductive agents and a huge volume expansion during the charge and discharge process. The high solubility of sulphur in common electrolytes causes instability in S cathodes and a loss of reversible capacity during cycling [178]. S cathodes, therefore, need to be combined with hard carbon as a functional host to prevent a sulphur shuttle effect and improve the reversibility of electrochemical reactions. Li et al [166] synthesized sodium sulphide (Na₂S) nanospheres embedded in hard carbons with a hollow structure and high conductivity, which shortens the ion diffusion pathway (figure 11(g)). As a result, a high initial discharge capacity of 980 mAh g⁻¹ at high current densities of 1.4 A g⁻¹ can be achieved (figure 11(h)). After 100 cycles, the reversible capacity can stabilize at 600 mAh g⁻¹. S-hard carbon composite cathodes can possess enhanced electrochemical reactivity by alleviating the volume change and shuttle effect of sulphur, resulting in excellent rate and cycling performance at various current densities.

Another kind of advanced cathode that can be employed for next-generation sodium batteries is the air cathode. The promising energy density of sodium-air batteries has widely motivated scientists to develop air cathodes. Sodium-air batteries utilize air from outside to participate in the reaction inside the batteries instead of the self-contained battery technologies mentioned above. High overpotential and a limited rate performance are two main restrictions of air cathodes, which result from the sluggish kinetics of oxygen reduction and evolution reactions (ORR/OER) [179]. Hard carbons with specific structures can be used as an effective bifunctional catalyst to boost the kinetics of electrochemical reactions inside the battery at the cathode side. Moreover, hard carbons, as functional hosts, are also able to hold solid discharge products like NaO₂ at the same time and provide smooth pathways for mass transport, including air and electrolytes [180]. Zhang et al [181] prepared a co-embedded N-doped hard carbon fabric as a free-standing functional host for an air cathode. A superior electrochemical performance containing a high discharge capacity, a low charge overpotential and a stable cycle life (∼112 cycles) all proved the positive effects of hard carbons in improving ORR/OER reactivity and mass transport. The successful development of next-generation high-performance air cathodes cannot happen without hard carbon functional hosts.