Johnson Matthey Technol. Rev., 2022, 66, (1), 44
Sodium-Ion Batteries: Current Understanding of the Sodium Storage Mechanism in Hard Carbons
Optimising properties to speed commercialisation
- Jack R. Fitzpatrick, Sara I. R. Costa, Nuria Tapia-Ruiz*
- Department of Chemistry, Lancaster University, Lancaster LA1 4YB, UK; The Faraday Institution, Quad One, Harwell Science and Innovation Campus, OX11 0RA, UK
Received 14th April 2021; Revised 29th June 2021; Accepted 29th June 2021; Online 29th June 2021
In recent years, sodium-ion batteries (NIBs) have been explored as an alternative technology to lithium-ion batteries (LIBs) due to their cost-effectiveness and promise in mitigating the energy crisis we currently face. Similarities between both battery systems have enabled fast development of NIBs, however, their full commercialisation has been delayed due to the lack of an appropriate anode material. Hard carbons (HCs) arise as one of the most promising materials and are already used in the first generation of commercial NIBs. Although promising, HCs exhibit lower performance compared to commercial graphite used as an anode in LIBs in terms of reversible specific capacity, operating voltage, initial coulombic efficiency and cycling stability. Nevertheless, these properties vary greatly depending on the HC in question, for example surface area, porosity, degree of graphitisation and defect amount, which in turn are dependent on the synthesis method and precursor used. Optimisation of these properties will bring forward the widespread commercialisation of NIBs at a competitive level with current LIBs. This review aims to provide a brief overview of the current understanding of the underlying reaction mechanisms occurring in the state-of-the-art HC anode material as well as their structure-property interdependence. We expect to bring new insights into the engineering of HC materials to achieve optimal, or at least, comparable electrochemical performance to that of graphite in LIBs.
At present, it is evident that a significant reduction or, better yet, a complete eradication in the use of fossil fuels is required towards achieving a more sustainable future, and 195 countries have signed the Paris agreement as of January 2021 towards limiting the global temperature increase to less than 2°C this century (1). This commitment will be achieved by transitioning to renewable energy generation and phasing out petrol and diesel vehicles in favour of electric vehicles (EVs). In this scenario, battery technology will have a major role to tackle the energy crisis and environmental problems. To date, LIBs are the most widely used battery technology due to their superior gravimetric and volumetric energy densities and cycle life (2, 3). However, the increasing demand for energy storage devices has raised concerns regarding their sustainability and cost-effectiveness, leading to the resurgence of NIB research in recent decades (4–11).
Sodium is significantly more abundant and uniformly distributed in the Earth’s crust compared to the scarcer and unevenly distributed lithium (Table I) (11–16). Consequently, the cost of sodium salts is much lower compared to lithium salts (for example, US$154 tonne–1 for sodium carbonate vs. US$8000 tonne–1 for lithium carbonate in 2020) (12). Furthermore, the working principle of both systems is similar (see Figure 1), and therefore, NIBs are considered a drop-in technology that can benefit from existing LIB manufacturing facilities. Another advantage of NIBs is the possibility to use low-cost and sustainable elements in their cathode materials (for example, iron, manganese, magnesium and titanium), moving away from cobalt, which is widely used in LIB cathodes (for example, nickel manganese cobalt (NMC) materials) (17, 18). However, the overall lower cost of NIBs can be mainly attributed to the use of aluminium (US$0.3 m–2) as the current collector for the negative electrode (anode) compared to the more expensive copper (US$1.2 m–2) used in LIBs (17). Unlike lithium, sodium does not alloy with aluminium at low potentials, allowing its use at both electrodes (19). Furthermore, since aluminium is lighter than copper, its use boosts the total energy density of the NIB.
Some companies such as Faradion (UK), Hi-Na (China) and Tiamat (France), have already demonstrated the feasibility of NIB technology (20). For instance, in 2018, Faradion’s 12 Ah NIB full cells exhibited a maximum energy density of 140 Wh kg–1, similar to the typical range of energy densities currently offered by commercially available lithium iron phosphate (LFP) // graphite systems (140–175 Wh kg–1). Additionally, it was predicted that by the end of 2020, 160 Wh kg–1 would be achieved in 32 Ah full cells, with cycle life matching that offered by some LFP // graphite cell suppliers (~3000 cycles) (20). Furthermore, current NIBs offer a 25–30% cost reduction in terms of US$ kWh–1 compared to current LFP // graphite systems (20). Today’s commercial NIBs are already beginning to replace lead-acid batteries in low energy density applications, such as starter-lighting-ignition batteries in internal combustion engine vehicles and backup energy storage applications. In the near future, NIBs will likely compete with LFP // graphite LIBs for mid-range energy density applications, that require a high degree of safety or rate performance, such as electric scooters, buses, power tools and stationary energy storage systems supplying the grid (14, 20). However, current commercial NIBs cannot compete in terms of energy density with the state of the art NMC/lithium nickel cobalt aluminium oxide (NCA) // graphite lithium-ion systems in high energy density applications, such as long-range EVs. This is due, in part, to the absence of high-performance anodes that can compete with the current graphite-based anodes in LIBs. Despite this, the development of NIBs has been much faster compared to that of commercial LIBs since their conception (20). Therefore, NIBs may become competitive with high energy density LIB systems very soon.
Although the raw materials that make up a NIB are more cost-effective than the equivalent ones for LIBs, it is their cost as a function of energy density that truly matters (17, 21, 22). Hence, successful commercialisation of NIBs will depend on how much improvement is attained in terms of energy density, long-term stability and power density, while keeping production costs low. This will be partly achieved through the development of new materials and the optimisation of existing ones. Currently researched anode materials can be categorised into three groups based on their sodium storage mechanism: (a) intercalation/insertion (23–25); (b) alloying (26–29); and (c) conversion (26, 30). Figure 2 shows examples of materials in each group and their theoretical energy density when used in a full cell. The most used materials in today’s NIBs are those in the first category, which can achieve specific capacities within the range of 100–400 mAh g–1. Examples include carbonaceous materials such as expanded graphite (31) and non-graphitisable carbon/HC, as well as titanium-based materials such as titanium oxide (TiO2) and titanates (for example, Na2O·nTiO2) (32, 33, 34).
When used in NIB systems, graphite, the most commonly used anode in commercial LIBs, is unable to intercalate a significant amount of Na+ ions with commercial carbonate-based electrolytes, achieving a specific capacity of ~35 mAh g–1 (35, 36). However, disordered carbons are able to reversibly store significant amounts of Na+ ions, especially HCs. However, there is no consensus on how Na+ ions are stored during the first discharge (i.e. sodiation process). This appears to be a result of the diverse structural properties exhibited by HCs when synthesised using different methods and precursors. In this review, we will summarise the current understanding of the sodium storage mechanism occurring in HCs, and draw correlations between the intrinsic structural properties found in these materials and their effect on the storage mechanism, as a means to advance in their optimisation. Finally, we provide a summary of the main challenges and possible directions towards the future optimisation of HCs.
In NIBs, graphite intercalates a negligible amount of Na+ ions when using commercial carbonate ester-based electrolytes, with specific capacities as low as 35 mAh g–1 (NaC64) (35, 37). Theoretical calculations show that the intercalation of Na+ ions into graphite to form graphite intercalation compounds (GICs) NaC6/NaC8, is thermodynamically unfavourable, unlike lithium insertion to form LiC6 (38–40). Instead, it is more favourable for the Na+ ions to be deposited on the surface of the graphite as metallic sodium (37). This is often attributed to the larger ionic radius of Na+ vs. Li+ (Table I), which prevents sodium intercalation between the graphitic layers. Yet, sodium is the exception among all alkali metals, as MC6 and MC8 GICs are readily formed with heavier Group I elements (for example, potassium, rubidium and caesium) (41). Thus, it is clear that the ionic radius is at least not the only accountable factor for the unsatisfactory electrochemical performance of graphite in carbonate-based electrolytes. However, the possibility of the use of graphite as an anode in NIB systems has been reignited in recent years, as in 2014 Jache et al. reported the reversible co-intercalation of Na+ ions into graphite through the use of diethylene glycol dimethyl ether (diglyme) electrolyte solvent (42). A reversible capacity of ~100 mAh g–1 was achieved with a high initial coulombic efficiency (ICE) of 90%, excellent cycling stability over 1000 cycles and high rate capability, maintaining ~75 mAh g–1 at 1 C. This finding sparked further work exploring the Na+ co-intercalation mechanism at graphite (43–46). However, most of the capacity resides as part of a high voltage plateau in the discharge/charge profile between ~0.8–0.6 V (vs. Na+/Na), leading to low energy density values.
On the other hand, HCs are the material of choice among all anode materials currently tested in NIBs. They are already used in the first generation of commercial NIBs (20), despite exhibiting lower specific capacity compared to graphite in commercial LIBs (200–300 vs. ~372 mAh g–1 respectively) (47–49). All the synthetic routes developed to produce HCs share in common a core pyrolysis step (i.e. high-temperature carbonisation using temperatures from 800°C to 2000°C) of an oxygen-rich organic precursor under an inert atmosphere (for example, argon or nitrogen) (50, 51). Although precursors can be synthetic, research mainly focuses on sustainable plant-derived precursors such as glucose, sucrose, cellulose and lignin (48, 52, 53). HC precursors cannot be converted into graphite, even at extremely high temperatures (>3000°C), and therefore, are often termed as non-graphitisable carbons (50, 52–54).
Structurally, HCs differ significantly from graphite, i.e. HCs exhibit a disordered structure, with short-range crystalline domains (Figure 3(a)) (55–59), while graphite exhibits long-range ordered stacking of graphene layers with a well-defined interlayer spacing (3.3 Å) (25). However, as the temperature of synthesis increases, HCs become structurally more ordered, with increasingly large graphitic domains (Figure 3(a)). Franklin first proposed a structural model for HCs (59), followed by Dahn et al. with the so-called ‘house of cards’ or ‘falling cards’ model (56–58). According to this model, HCs consist of randomly oriented graphitic domains (two or three stacked layers) with expanded interlayer spacing (3.6–4.0 Å), and different curvatures that are interconnected by highly disordered regions. This mismatch of ordered and disordered domains causes the presence of closed pores (dotted red lines in Figure 3(b)), between the randomly orientated graphitic crystallites (56–58). Also, HCs tend to contain a higher content of heteroatoms (mainly oxygenated groups) than graphite, providing more defects and sites for Na+ ion adsorption (52).
Although the typical specific capacity values achieved by HCs are significantly lower compared to alloying- and conversion-based anodes (300 mAh g–1 vs. 2600 mAh g–1 for phosphorus, for example) (Figure 2), they exhibit superior cycling stability due to minimal volume changes during cycling (~1–2.5%) (60). As such, HC containing cells can exhibit lifetimes up to 1000 cycles compared to 50 and 20 cycles for antimony and lead alloying materials, respectively, which show large volume changes in the order of ~300% during cycling (24, 26, 27). Additionally, the working potential of HC (0.1 V vs. Na+/Na) is one of the lowest among intercalation anode materials (Figure 2), which is a great benefit for attaining high energy density (61, 62). Lastly, they are cost-effective, enabling the production of batteries at a lower cost.
Nevertheless, HC materials face various challenges that need to be addressed before achieving their full commercialisation. To date, the sodium storage mechanism remains elusive, as many research groups have reported contradictory studies in the last years. Attaining this understanding is paramount to developing feasible strategies to accomplish superior electrochemical performance. The complex structure of HCs implies that there are multiple sodium storage pathways, and different synthesis methods or precursors result in different structural features. As a result, significant variations in the electrochemical behaviour are observed, making it difficult to establish a universal storage mechanism model. Therefore, it is reasonable to consider that multiple storage mechanisms exist, depending on the (micro)structure and properties of the as-synthesised HC.
Furthermore, cycling typically shows a low ICE induced by a large surface area exposed to the electrolyte (usually below 80%, (47, 48, 63) which contrasts to the 85–95% ICE values reported for graphite (63, 64)). This leads to an extensive reduction of the electrolyte during the first discharge, with subsequent formation of a solid electrolyte interface (SEI) layer, and partially irreversible trapping of Na+ ions in the structure (65–68). The composition of the SEI formed on graphite in LIBs has been extensively studied and optimised in such a way that parameters including cycling stability, rate capability and safety have greatly improved (69–72). However, similar studies in NIBs are still in their infancy, partly due to the higher instability of the SEI components in NIBs (73–77).
Sodium Storage Mechanism
Extensive research efforts have been made in the last two decades to provide insight into the sodium storage mechanism in HCs. Many studies have reported that the structural and microstructural features of HCs, including pore size, pore volume and terminal groups, are influenced by the synthetic conditions (for example, carbonisation temperature, dwelling time and atmosphere) and precursors choice. This, in turn, has a direct repercussion on several parameters that define the electrochemical behaviour of HCs, including specific capacity, cycling stability, coulombic efficiency, rate capability, wettability, and ultimately, the storage pathway of Na+ ions. For reference, a typical galvanostatic discharge/charge voltage profile of HC at low current rates (for example, 0.1 C and 0.2 C) is shown in (Figure 4(a)). This consists of two distinct features: a sloping region at a potential above ~ 0.1 V vs. Na+/Na, and a plateau region below ~ 0.1 V (78). Given the complex structural nature of HCs, researchers concur with the existence of a multi-step storage mechanism upon cycling which consists of the following reactions (52, 53): (a) intercalation of Na+ ions between the pseudo-graphitic layers; (b) adsorption of Na+ ions at reactive surfaces and defect sites (for example, vacancies and dangling bonds at the edges of the pseudo-graphitic domains (78)); and (c) filling of closed nanopores with Na+ ions (see Figure 4(b)). Overall, five different mechanistic models have been proposed in the literature (Figure 4(c)–4(g)) (78–84).
Dahn et al. were the first to report a sodium storage model for HCs, which was based on an intercalation-pore filling mechanism (Figure 4(c)) (37, 79, 85). The sloping region was assigned to the insertion of Na+ ions into the interlayer spacing of pseudo-graphitic domains, and the plateau region to the insertion/adsorption of sodium into closed nanopores. Observations were supported by in situ wide and small-angle X-ray scattering data, wide-angle X-ray scattering (WAXS) and small angle X-ray spectroscopy (SAXS), respectively (37). WAXS data showed a progressive decrease in intensity of the (002) reflection (attributed to the interlayer spacing of the graphitic layers) in the sloping potential region during discharge. This is consistent with the introduction of Na+ ions between the layers. However, no accompanying shift in 2θ values was seen for the (002) reflection. This process was found to be partially reversible upon desodiation. Furthermore, SAXS data showed a reversible decrease in intensity of scattering in the plateau region, suggesting the presence of scattering species (Na+ ions) entering the closed nanopores. From this, they were able to calculate the difference in electron density between the surrounding carbon matrix and the closed pores. A partially reversible decrease in electron density contrast upon sodiation across the plateau region was observed, which was attributed to the filling of closed nanopores with sodium (37).
This model was further supported by Komaba et al. who used ex situ X-ray diffraction (XRD) to show a gradual expansion of the (002) interlayer spacing in the sloping region during sodiation, demonstrating the occurrence of an intercalation process in this region (86). They also monitored changes in electron density with ex situ SAXS upon sodiation, corroborating previous results from Dahn et al. for closed nanopore filling in the plateau region (37, 85). More recently, in situ electrochemical dilatometry (ECD) studies have shown that there is a non-linear expansion of HC during sodiation, which increases in the sloping region and then levels off in the plateau region (60). Furthermore, it has been shown that the interlayer spacing in HCs decreases with increasing carbonisation temperature (82, 87–90), which is followed by a decrease in sloping capacity. Overall, these findings provide evidence for an intercalation process occurring in the sloping region and pore filling in the plateau region in accordance with the original model proposed by Dahn et al (37, 85).
However, as HCs have become extensively investigated in recent years, alternative storage models have been proposed as a result of conflicting experimental data (Figure 4(d)–4(g)). As mentioned earlier, differences arise from the different choices of precursors or synthesis conditions which strongly influence the properties of HCs and their respective electrochemistry (82, 87, 89–92). For instance, some authors have observed a shift in the (002) reflection solely in the plateau region from ex situ XRD data (67, 78, 91, 93) and in situ XRD data (81), in discordance with data reported by Stevens and Dahn (37, 85) and Komaba et al. (86), while others have not observed any changes in the lattice spacing over the whole sodiation process from in situ XRD and transmission electron microscopy (TEM) data (82, 83).
It has been widely observed that the surface area and defect site concentration of HCs decrease with increasing carbonisation temperature (Figure 5(a)) (78, 89, 92–94). Many authors have also observed a simultaneous decrease in the specific capacity of the sloping region (Figure 5(b)), with many reports showing a positive correlation between a decrease in sloping capacity and a decrease in surface area and defect site concentration. This link between the surface area, defect site concentration and sloping capacity, with increasing carbonisation temperature, suggests an adsorption process of Na+ ions occurring at the open pores and defect sites present at the surface of HCs, as opposed to an intercalation process during the sloping region (Figure 4(d)–4(f)) (78, 89, 92, 94). Bommier et al. suggested that different defect sites may have their own Na+ ion binding energy, and therefore, sodiation potential, resulting in a sloping shape in the galvanostatic profile (78). However, it is noteworthy to point out that HCs with higher surface and defects may trigger more side reactions with the electrolyte, increasing the specific capacity along the sloping region during the initial cycles due to SEI formation, in addition to the pseudocapacitive contribution of the adsorption process (54, 90).
The surface area, pore sizes or pore volume of HCs reported in the literature are typically derived from Brunauer, Emmett and Teller (BET) analysis of N2/CO2 adsorption data. It is worth noting that BET derived pore information does not consider internal closed pores present within the bulk of HC particles (88, 95). The presence of defects and their concentration in HCs have been assessed with Raman spectroscopy. For example, Sun et al. calculated the disorder level and defect concentration in HCs synthesised at various carbonisation temperatures by assessing the intensity ratio between the G and D bands (IG/ID), where IG is the intensity of the band from planar sp2 carbon atoms of perfect crystalline graphite, and ID represents the intensity of the defect-induced peak related to Brillouin zone-edge phonons (K point, sp3 carbon) (Figure 6(a) and 6(b)) (89). They observed an increase in the IG:ID ratio, ranging from 0.38 to 9.19 with increasing carbonisation temperature (from 600°C to 2500°C), indicative of a decrease in defect concentration and degree of disorder (96). In addition, they observed that the specific sloping capacity increases linearly with the ID:(IG+ID) ratio, further supporting an adsorption process at the surface occurring within the sloping region (Figure 6(b)). Also, other reports have shown a decrease in the sloping capacity with decreasing heteroatom content (oxygen and nitrogen containing groups), which normally occurs at higher carbonisation temperature, providing further evidence of an adsorption mechanism in the sloping region (82, 88, 92).
Studies have also attempted to understand the factors influencing the plateau capacity, and respective storage mechanism, with the influence of carbonisation temperature, also being widely investigated (82, 88, 89, 92, 94). As can be seen in Figure 5(b), the plateau capacity increases with carbonisation temperature. Nonetheless, this trend is observed only up to certain temperatures (1200–1400°C), depending on precursors or synthesis conditions, beyond which the capacity of the plateau region starts decreasing (89, 92, 94). Gomez-Martin et al. found that the calculated d002 interlayer spacing (from pair distribution function (PDF) data) of pristine olive-stone derived HCs at various carbonisation temperatures followed this trend, increasing up to 1400°C, after which it started decreasing (92). This allowed them to show a linear relationship between the d002 spacing and the plateau capacity (Figure 7). It is worth noting that not all studies exhibit this same trend in d002 values, with it being more common that a simple decrease in d002 with increasing carbonisation temperature is observed (60, 68, 88, 94, 97–99). However, it is widely observed that there is a strong correlation between the contribution of the plateau capacity to the calculated average width (La) and thickness (LC) (calculated using the d002 reflection from the XRD data) of the graphite-like domains. Therefore, suggesting that as the degree of graphitisation increases, the graphite crystallites grow and provide more sites for the storage of Na+, thus providing evidence for an intercalation process occurring during the plateau region (Figure 4(d) and 4(f)). In addition, Qiu et al. reported further evidence using in situ XRD for reversible intercalation within the plateau region (81). They showed that the (002) broad reflection at 26.6° 2θ gradually splits to form a new sharp reflection at 26.2° 2θ (larger interlayer spacing) near the start of the first discharge plateau which shows the preference of Na+ to intercalate in graphitic domains with larger spacing.
If we now consider the pore-filling mechanism and its possible contribution to the plateau capacity (Figure 4(c), 4(e), 4(f) and 4(g)). It is the closed pores, enclosed and isolated within the bulk of the HC structure, that are of interest to the pore-filling mechanism (60, 88, 92, 100). Hu et al. showed evidence for a pore-filling mechanism in the plateau region using X-ray photoelectron spectroscopy (XPS) and TEM data (83). Their ex situ Na 1s XPS data from HC samples after etching (60 nm) (Figure 8(a)), showed a shift of the sodium binding energy towards higher energies, close to that of metallic sodium (~1072 eV) upon discharge below 0.12 V vs. Na+/Na, which is consistent with deposition of metallic sodium in the closed pores. Additionally, TEM images (Figure 8(b)) showed blurring of the closed nanopores within the bulk structure and at the edges of the graphitic layers in the HC discharged to 0 V vs. Na+/Na, which the authors attributed to sodium deposition in these regions. In the same year, Grey et al. published a study supporting this model with 23Na magic angle spinning nuclear magnetic resonance (MAS-NMR) (Figure 9) (101). It was observed that the resonance peak at 0 ppm shifted to approximately 760 ppm near to that of metallic sodium in the low-voltage plateau, consistent with the formation of metallic sodium in the closed pores. Furthermore, PDF analysis corroborated these findings, showing an extra phase with atom-atom distances of ~10 Å appearing during the plateau region (0.05 V), similar to those observed in sodium metal. Furthermore, Hu et al. used SAXS to calculate the closed pore volume of HCs produced at different temperatures (88). They found that the volume of closed pores increased from 0.11 cm3 g–1 to 0.29 cm3 g–1 with increasing carbonisation temperature (from 1200°C to 1600°C) which led to an enhancement of the plateau capacity.
The models described so far support the presence of a two-stage storage mechanism. More recently, a fourth type of model that combines the three postulated sodium storage pathways has been proposed by multiple authors (Figure 4(f) and 4(g)) (78, 84, 94, 102, 103). The first type of three-stage model (Figure 4(f)) was initially proposed by Bommier et al. who performed galvanostatic intermittent titration (GITT) measurements on HC to determine the Na+ ion diffusivity at different potentials (Figure 10) (78). Data showed higher Na+ ion diffusivity values in the slope compared to the plateau region, suggesting that intercalation, being a more kinetically impeded process, occurs at low potentials while adsorption occurs at higher potentials. However, at a potential below 0.05 V (end of the plateau), the diffusion increased again, consistent with the deposition of near-metallic sodium into the closed nanopores. This trend in Na+ ion diffusivity has been reported in other recent works (91, 94). These observations also explain the finding that more disordered HCs (produced at lower carbonisation temperatures) which exhibit greater defect concentrations and larger sloping capacities, provide superior rate capability, with the plateau capacity decreasing significantly more than the sloping capacity at high current rates (Figure 11(a)–11(b)) (89, 94). Figure 12(a)–12(c) shows an example of the long-term cycling stability offered by a range of sycamore fruit seed derived HCs, carbonised at various temperatures (900–1500°C) (97). The cycling stability decreases with increasing carbonisation temperature (Figure 12(a)), being optimal at 1100°C (at a 0.4 C rate). The decrease in cycling stability is attributed to a significant decrease in plateau capacity upon (de)sodiation cycles (Figure 12(c)). The authors explain the decrease in stability by the smaller interlayer spacing observed for HCs produced at higher temperatures, causing proportionally larger volume changes during Na+ intercalation that destroy the graphitic domains. In contrast, the sloping capacity remains stable for all of the HCs over 300 cycles (Figure 12(b)). Puravankara et al. observed similar behaviour in their sucrose derived HCs carbonised at 1000°C (104). The plateau capacity decreased from 215 mAh g–1 to ~150 mAh g–1 over 50 cycles when cycling at 0.1 C rate (30 mAh g–1), while the sloping capacity remained almost constant at 170 mAh g–1. In addition, increasing the current rate above 1 C caused the contribution from the plateau region to drop substantially.
An alternative three-stage storage mechanism has recently been proposed by Au et al. (here named three-stage model 2) (Figure 4(g)) (84). In this model, the capacity obtained in the sloping region has a contribution from both intercalation and adsorption processes between 2.5 V and 0.1 V. As the potential further decreases and these sites are filled, they suggest that pore filling occurs with the formation of quasi-metallic sodium. Their evidence relied on DFT calculations consistent with sodium binding energy in bilayer and monolayer graphitic models. These showed that, during the sloping region, sodium storage at surface defects initially dominates. Once these become saturated, sodium intercalation within graphitic bilayers becomes energetically favourable too. Moreover, DFT calculations determined that the maximum theoretical capacity derived from defect adsorption only accounts for 48–79% of the total specific capacity in the sloping region. They also observed that with increasing carbonisation temperature (1000–1900°C) the O:C ratio, degree of defects and interlayer spacing all decrease concurrently with a decrease in sloping capacity. Therefore, they suggest two different storage mechanisms exist in the sloping region, contrary to many models previously reported (Figure 4(d)–4(f)).
The theoretical calculations were complimented with experimental analysis including ex situ 23Na MAS-NMR measurements of HC samples carbonised at varying temperatures discharged to 5 mV. They observed a peak consistent with the presence of quasi-metallic sodium at 750 ppm for HC samples which also exhibited a plateau in their load curves, consistent with the findings of Grey et al. (101). For the sample treated at 1000°C which only exhibits a sloping region, this peak is not present. This implies that sodium stored in the sloping region is mainly ionic in nature, consistent with sodium adsorbed at defects, between interlayers and at open pore surfaces. Also, with increased temperature and increased closed pore diameter (determined via SANS), this quasi-metallic peak grows in intensity and shifts to higher ppm towards that of metallic sodium (1135 ppm). This growing metallic character of the sodium with increasing closed pore size is consistent with the growth of sodium clusters within the closed pores during the plateau region.
At present, there is still no consensus regarding the sodium storage mechanism in HCs, especially in the low voltage plateau region. Strict comparisons between HCs cannot be made due to their rich and diverse nature, which lead to significant differences in electrochemical behaviour and sodium storage mechanism. In short, the lack of a HC benchmark to perform mechanistic studies is largely responsible for the disagreement among experimental findings. Furthermore, reliance on ex situ measurements to probe changes in the HC (micro)structure during cycling contributes to the current controversy. Hence, it is clear that one model will never fit all HCs and so it is reasonable to consider the existence of different models for different families of HCs. Thus, moving forward, optimisation strategies should rely on an in-depth understanding of individual families of HCs with shared properties, for example, HCs produced from identical precursors or similar carbonisation temperatures. Full understanding of these mechanisms will demand the use of operando and in situ advanced techniques to provide complementary data at different length scales. For instance, operando Raman spectroscopy has shown to be more reliable than ex situ studies, as the electrode washing process and air exposure causes sodium deintercalation from HC, which results in the G- and D-peaks reverting to their original states (105).
Continuing to develop a deep insight into the relationship between the (micro)structural properties of HCs and their electrochemical performance is crucial. This will allow optimisation of individual HCs tailored to specific applications requiring, for example, enhanced energy density, cycling stability or rate performance. We believe it is possible to further enhance the electrochemical performance of HCs using our current understanding of their sodium storage mechanism. In this regard, targets for future research should include:
(a) Increasing the specific capacity of HCs at the plateau region by:
(i) Expanding the interlayer spacing of HCs to facilitate Na+ ion intercalation, while still maintaining a high degree of order within the structure. This could be achieved through the introduction of dopants or pillaring agents into the structure. For example, 0.1 wt% sulfur and 3.0 wt% phosphorus-doped HCs showed larger interlayer spacings compared to an undoped sample (increasing from 3.77 Å to 3.83 Å and 3.95 Å respectively). This led to an increase in the first discharge plateau capacity of 28 mAh g–1 and 53 mAh g–1 for the sulfur and phosphorus doped HCs respectively, with respect to the undoped analogue (106).
(ii) Synthesising HCs with an increased number of closed nanopores. One way this can be achieved is through the introduction of pore-forming additives to HC precursors. One recent example used a magnesium oxide templating technique, which dispersed magnesium oxide nanoparticles throughout the glucose precursor. These acted as a template to form nanosized pores in the resultant hard carbon after high-temperature carbonisation. The formed HC exhibited an extremely high plateau capacity of 401 mAh g–1 (107).
(b) Increasing the ICE (100, 108). This is particularly vital for HCs tailored towards power applications since their storage capability is mostly achieved through the enhancement of the sloping capacity. Strategies to improve the ICE include:
(i) Producing low-porosity and low-defect HCs using high carbonisation temperatures (≤1200°C) to minimise adsorption processes occurring at the surface. However, it should be noted that this will decrease the sloping capacity, which will be detrimental to the power density. Therefore, it is always a balancing act and the target application (energy or power) of the HC should always be kept in mind during any optimisation process.
(ii) Using surface engineering strategies to minimise the irreversible surface reactions with Na+ ions. For example, by decreasing oxygen functionality using a reducing carbonisation atmosphere (109) or coating the HC particle surface with a defect-free or less defective material such as a layer of soft carbon (110).
(iii) Electrolyte optimisation (for example, choice of sodium salt, solvents and additives) with the aim to produce a stable and thin SEI layer at the HC surface (111).
(iv) Pre-sodiation of the HC anode in full-cells to compensate for the sodium loss in the first cycle. This can be achieved by HC pre-sodiation in half cells before full cell assembly (112), or by cathode pre-sodiation through the introduction of sodium-containing sacrificial additives such as Na3P and NaN3 (113, 114).
(c) Using three-electrode cells and testing of HCs in full-cells to decouple sodium plating reactions from sodium intercalation and pore-filling processes occurring at the plateau region. Likewise, increasing the potential at which the plateau region occurs should be considered to minimise plating issues.
(d) Sustainability of the production process and precursor costs, with the price of HCs needing to be competitive with the current price for battery grade graphite (~US$20 kg–1).
However, our current understanding of HCs is still insufficient for their widespread commercialisation. We predict that the development of characterisation techniques adequate to investigate the properties of HCs will be crucial to comprehend their behaviour and improve their performance even further, towards their widespread commercialisation and the maturation of NIB technology.
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Jack Fitzpatrick is a PhD student and part of the Tapia-Ruiz group in the Department of Chemistry at Lancaster University, UK. His project involves studying hard carbon material for use as a sodium-ion battery anode, with a focus on the characteristics of the SEI that forms on their surface. He previously obtained an MChem from the University of York, UK, in 2020.
Sara Costa is currently a Senior Research Associate at Lancaster University. She obtained her PhD from Loughborough University, UK, in 2020, with a thesis on anode materials for sodium-ion batteries. Her current research focuses on the characterisation and improvement of anode materials for both sodium-ion and lithium-ion batteries.
Nuria Tapia-Ruiz obtained her PhD from the University of Glasgow, UK, in 2013. She then worked as a Research Fellow in the team of Professor Bruce at the University of St Andrews and the University of Oxford, UK, in 2013–2016. Currently, she is a Senior Lecturer in the Department of Chemistry at Lancaster University. Her research interests include the understanding of the structure-property-performance relationships in materials for energy storage such as monovalent and multivalent batteries and supercapacitors.