In the pursuit of sustainable energy storage solutions, the development of efficient and cost-effective battery systems has become paramount. As a researcher deeply engaged in this field, I have focused on exploring alternative materials for sodium-ion batteries, which offer a promising avenue due to the abundance of sodium resources compared to lithium. The sodium-ion battery represents a key technology for large-scale applications, such as electric vehicles and grid storage, but its commercialization hinges on the availability of suitable anode materials. Hard carbon, derived from biomass precursors, has emerged as a viable candidate due to its disordered structure and ability to accommodate sodium ions. In this study, I investigated the regulation of pore structure in shaddock peel-based hard carbon using a template method, aiming to enhance its sodium storage performance for sodium-ion battery applications. The core innovation lies in employing magnesium oxide as a template, coupled with magnesium gluconate as a pore-forming agent, to create a hierarchical porous architecture that facilitates ion diffusion and storage. Through systematic characterization and electrochemical testing, I demonstrated that this tailored hard carbon anode delivers high capacity and stability, underscoring its potential for advancing sodium-ion battery technology.
The global energy landscape is rapidly evolving, driven by the need for renewable integration and electrification of transportation. Lithium-ion batteries have dominated the energy storage market, but concerns over lithium scarcity and cost have spurred interest in alternatives. The sodium-ion battery, leveraging the similar electrochemistry of sodium to lithium, presents a compelling solution. Sodium is Earth-abundant, geographically widespread, and inexpensive, making sodium-ion batteries ideal for scalable deployment. However, the larger ionic radius of sodium ($\text{Na}^+$, 1.02 Å) compared to lithium ($\text{Li}^+$, 0.76 Å) poses challenges in identifying anode materials that can reversibly intercalate sodium ions without significant volume expansion or capacity fading. Graphite, the standard anode in lithium-ion batteries, exhibits poor sodium storage due to thermodynamic incompatibility, with a limited interlayer spacing of approximately 0.335 nm. Hard carbon, characterized by its turbostratic structure comprising randomly oriented graphene sheets, defects, and micropores, offers a larger interlayer distance (typically 0.37–0.40 nm) and abundant active sites, making it suitable for sodium-ion battery anodes. The sodium storage mechanism in hard carbon involves a combination of adsorption on defect sites, intercalation between graphene layers, and pore filling, which collectively contribute to its capacity. Specifically, the discharge curve of a sodium-ion battery with a hard carbon anode typically shows a slope region above 0.1 V (attributed to surface adsorption and defect storage) and a plateau region below 0.1 V (associated with pore filling and intercalation). Optimizing this plateau capacity is crucial for achieving high energy density in sodium-ion batteries.
Biomass-derived hard carbons are particularly attractive due to their low cost, sustainability, and tunable properties. Shaddock peel, a common agricultural waste rich in cellulose and hemicellulose, serves as an excellent carbon precursor because it yields high carbonization efficiency and inherently contains heteroatoms like oxygen and nitrogen that can enhance conductivity and surface reactivity. However, the inherent pore structure of biomass-derived carbons is often irregular and limited, necessitating post-synthesis modifications to improve sodium ion accessibility. Pore engineering via templating methods has proven effective in creating controlled porosity, where sacrificial templates generate pores upon removal. In this work, I utilized magnesium oxide (MgO) as a hard template, given its stability at high temperatures and ease of removal through acid washing or reduction. Magnesium gluconate was incorporated as a co-precursor because it decomposes to form MgO nanoparticles during pre-carbonization, acting as both a template and a pore-forming agent. By varying the mass ratio of shaddock peel to magnesium gluconate and employing different mixing techniques (dry powder mixing vs. solution mixing), I aimed to control the dispersion of MgO and, consequently, the pore size distribution in the final hard carbon. The hypothesis was that a well-distributed porous network would shorten ion diffusion paths, increase electrode-electrolyte contact area, and provide more storage sites, thereby boosting the performance of the sodium-ion battery.
The experimental procedure began with the collection and preparation of shaddock peel. Fresh shaddock peels were washed, dried at 80°C for 24 hours, and ground into a fine powder using a ball mill. Magnesium gluconate (C12H22MgO14) was procured as a reagent-grade chemical. For the template synthesis, I prepared two series of composites by varying the mass ratio of shaddock peel to magnesium gluconate: 3:1, 2:1, 1:1, 1:2, and 1:3. Two mixing methods were employed to assess the effect of precursor dispersion. In the dry mixing method, denoted as “dry method,” the shaddock peel powder and magnesium gluconate were mechanically blended in an agate mortar for 30 minutes. In the wet mixing method, denoted as “wet method,” magnesium gluconate was first dissolved in deionized water under magnetic stirring, followed by the addition of shaddock peel powder to form a slurry; this was then stirred for 2 hours, filtered, and dried at 100°C overnight. The mixed precursors were subjected to pre-carbonization in a tubular furnace under argon atmosphere at 600°C for 2 hours with a heating rate of 5°C/min. This step decomposed the organic components and converted magnesium gluconate into MgO nanoparticles embedded in a carbon matrix. The pre-carbonized products were then treated with 2 M hydrochloric acid (HCl) for 12 hours to remove most of the MgO, washed with deionized water until neutral pH, and dried at 100°C. Finally, high-temperature carbonization was performed in an induction furnace at 1400°C for 2 hours under argon flow (10°C/min heating rate) to graphitize the carbon and further eliminate residual MgO via reduction and evaporation. The obtained hard carbon samples were labeled accordingly, e.g., HC-wet 1:2 for wet-mixed sample with 1:2 ratio, and HC-dry 1:2 for dry-mixed counterpart. For comparison, a pristine shaddock peel-derived hard carbon (labeled as HC-pristine) was prepared without any additives under identical carbonization conditions.
Material characterization involved multiple techniques to elucidate the structural and chemical properties. X-ray diffraction (XRD) patterns were recorded using a Bruker D8 Advance diffractometer with Cu Kα radiation (λ = 1.5406 Å) over a 2θ range of 10–80°. The interlayer spacing (d002) was calculated using the Bragg equation: $$2d\sin\theta = n\lambda$$ where θ is the diffraction angle, λ is the wavelength, and n=1 for the (002) peak. Scanning electron microscopy (SEM) images were obtained with a JEOL JSM-7800F microscope to observe morphology and pore distribution. Fourier-transform infrared (FT-IR) spectroscopy was performed on a Nicolet iS50 spectrometer in the range of 400–4000 cm−1 to identify functional groups. X-ray photoelectron spectroscopy (XPS) analysis was conducted on a Thermo Scientific K-Alpha system to probe surface composition and bonding. Nitrogen adsorption-desorption isotherms were measured at 77 K using a Quantachrome Autosorb-iQ analyzer; specific surface area was calculated via the Brunauer-Emmett-Teller (BET) method, and pore size distribution was derived from the density functional theory (DFT) model. The total pore volume (Vtotal) and average pore diameter (Davg) were also determined. Electrochemical evaluations were carried out by assembling CR2025 coin cells in an argon-filled glovebox with oxygen and moisture levels below 0.01 ppm. The working electrode was fabricated by mixing the hard carbon active material, acetylene black conductive agent, and polyvinylidene fluoride (PVDF) binder in a weight ratio of 8:1:1 in N-methyl-2-pyrrolidone (NMP) solvent, coating the slurry on copper foil, and drying at 80°C under vacuum. Sodium metal foil served as the counter/reference electrode, and a glass fiber separator (Whatman GF/D) was soaked with electrolyte consisting of 1 M sodium hexafluorophosphate (NaPF6) in diethylene glycol dimethyl ether (DIGLYME). Cyclic voltammetry (CV) was performed on a CHI660E electrochemical workstation at scan rates from 0.1 to 1.0 mV/s between 0.01 and 3.0 V vs. Na/Na+. Electrochemical impedance spectroscopy (EIS) measurements were conducted in the frequency range of 0.01 Hz to 100 kHz with an amplitude of 5 mV. Galvanostatic charge-discharge tests were executed on a Neware battery tester at various current densities from 30 to 1000 mA/g. The specific capacity (C) was calculated based on the mass of hard carbon using the formula: $$C = \frac{I \times t}{m}$$ where I is the current, t is the discharge time, and m is the active mass. All tests were performed at room temperature (25°C).
The structural analysis via XRD revealed distinct differences between the pristine and templated hard carbons. The XRD patterns of all samples exhibited two broad diffraction peaks around 23° and 44°, corresponding to the (002) and (101) planes of turbostratic carbon, respectively. For HC-pristine, the (002) peak was centered at 22.6°, yielding an interlayer spacing d002 of 0.38 nm as per the Bragg equation. In contrast, the templated samples showed a slight shift in the (002) peak position. For instance, HC-wet 1:2 had a peak at 24.1°, giving d002 = 0.37 nm, while HC-dry 1:2 showed a similar shift. This reduction in d002 indicates a slight graphitization enhancement due to the templating effect, yet all values remained larger than graphite’s 0.335 nm, facilitating sodium ion intercalation. The (101) peak was weakly defined, confirming the disordered nature of the hard carbons. The crystallite size along the c-axis (Lc) was estimated using the Scherrer equation: $$L_c = \frac{K\lambda}{\beta\cos\theta}$$ where K is the shape factor (0.89), λ is the X-ray wavelength, β is the full width at half maximum (FWHM) of the (002) peak, and θ is the Bragg angle. The calculated Lc values ranged from 1.2 to 1.8 nm, consistent with typical hard carbons. The incorporation of MgO template did not introduce impurity peaks, confirming its complete removal after acid washing and high-temperature treatment. To summarize the XRD data, I present a table comparing key parameters:
| Sample | (002) Peak Position (°) | d002 (nm) | FWHM (°) | Lc (nm) |
|---|---|---|---|---|
| HC-pristine | 22.6 | 0.38 | 5.2 | 1.5 |
| HC-wet 1:2 | 24.1 | 0.37 | 4.8 | 1.7 |
| HC-dry 1:2 | 23.8 | 0.37 | 5.0 | 1.6 |
| HC-wet 1:1 | 23.5 | 0.38 | 5.1 | 1.5 |
| HC-dry 1:1 | 23.3 | 0.38 | 5.3 | 1.4 |
Morphological observations through SEM provided visual evidence of pore formation. HC-pristine displayed irregular, sheet-like particles with a relatively smooth surface and no distinct porosity. In contrast, the templated samples exhibited abundant pores distributed throughout the carbon matrix. HC-wet 1:2 showed a homogeneous porous structure with pore sizes ranging from tens to hundreds of nanometers, attributed to the even dispersion of MgO nanoparticles via solution mixing. HC-dry 1:2 had a less uniform pore distribution, with some agglomerated regions, due to the inferior mixing of solid precursors. The porous architecture is crucial for electrolyte penetration and sodium ion transport, directly impacting the performance of the sodium-ion battery. Energy-dispersive X-ray spectroscopy (EDS) mapping confirmed the absence of magnesium after acid washing, verifying template removal. FT-IR spectra further supported the structural changes. All samples displayed characteristic absorption bands: a broad band at 3450 cm−1 (O–H stretching), peaks at 2810 cm−1 (C–H stretching in –CH2–), 2720 cm−1 (aldehyde C–H), 1600 cm−1 (C=C aromatic stretching), 1300 cm−1 (C–O stretching), and 775 cm−1 (C–H bending). The templated samples showed reduced intensity in these bands compared to HC-pristine, indicating a lower oxygen content due to higher carbonization degree, which is beneficial for electronic conductivity in sodium-ion battery anodes. No new peaks appeared, confirming that the wet mixing process did not induce chemical reactions beyond decomposition.
Nitrogen sorption analysis quantified the porous characteristics. HC-pristine had a low BET specific surface area of 2.526 m2/g, with negligible microporosity. The templated samples exhibited significantly enhanced surface areas and pore volumes. HC-wet 1:2 achieved a BET area of 8.012 m2/g, a total pore volume of 0.025 cm3/g, and an average pore diameter of 3.5 nm, indicating a mix of micropores and mesopores. HC-dry 1:2 had a lower BET area of 3.651 m2/g, volume of 0.012 cm3/g, and pore diameter of 4.2 nm. The wet method evidently produced a more developed porous network, as summarized in the table below:
| Sample | BET Surface Area (m2/g) | Total Pore Volume (cm3/g) | Average Pore Diameter (nm) | Micropore Volume (cm3/g) |
|---|---|---|---|---|
| HC-pristine | 2.526 | 0.005 | 6.2 | 0.001 |
| HC-wet 1:2 | 8.012 | 0.025 | 3.5 | 0.008 |
| HC-dry 1:2 | 3.651 | 0.012 | 4.2 | 0.003 |
| HC-wet 1:1 | 6.543 | 0.020 | 3.8 | 0.006 |
| HC-dry 1:1 | 4.210 | 0.015 | 4.0 | 0.004 |
The pore size distribution curves derived from DFT analysis revealed that HC-wet 1:2 had a prominent peak around 0.8 nm (micropores) and a broader peak around 3–5 nm (mesopores), whereas HC-dry 1:2 showed fewer micropores. This hierarchical porosity is advantageous for sodium ion storage: micropores provide active sites for adsorption, while mesopores facilitate ion diffusion. The formation mechanism can be described as follows: during pre-carbonization, magnesium gluconate decomposes into MgO nanoparticles that act as templates; upon acid washing, most MgO is leached out, leaving behind voids. During high-temperature carbonization at 1400°C, residual MgO is reduced by carbon to magnesium metal, which evaporates, creating additional pores. The overall pore volume (Vp) can be approximated by the equation: $$V_p = \frac{m_{\text{MgO}}}{\rho_{\text{MgO}}} \times \phi$$ where mMgO is the mass of MgO formed, ρMgO is its density (3.58 g/cm3), and φ is the removal efficiency. For a 1:2 ratio, the theoretical pore volume is around 0.028 cm3/g, close to the experimental value for HC-wet 1:2, confirming the template’s efficacy.
XPS analysis provided insights into surface composition and bonding. Survey spectra confirmed the presence of carbon (C 1s) and oxygen (O 1s), with trace nitrogen from the biomass. The C 1s peak was deconvoluted into components at 284.6 eV (C–C sp2), 285.2 eV (C–C sp3), 286.3 eV (C–O), and 288.9 eV (C=O). The O 1s peak showed contributions from C=O (531.5 eV) and C–O (533.0 eV). The atomic percentage of oxygen decreased from 12.5% in HC-pristine to 8.2% in HC-wet 1:2, indicating enhanced graphitization, which improves electronic conductivity for sodium-ion battery operation. The nitrogen content was below 1%, suggesting minimal doping effect. The surface functional groups can influence the solid electrolyte interphase (SEI) formation and initial coulombic efficiency. The higher carbonization temperature likely reduced surface oxygen groups, mitigating irreversible reactions with the electrolyte.
Electrochemical performance evaluation began with cyclic voltammetry (CV). The CV curves of HC-wet 1:2 at a scan rate of 0.1 mV/s exhibited reduction peaks at around 0.1 V and 0.01 V during the first cathodic scan, corresponding to SEI formation and sodium ion insertion into pores, respectively. In subsequent cycles, the curves became highly reversible, with oxidation peaks near 0.2 V indicating sodium extraction. The integral area of the CV curves reflects the capacity, and HC-wet 1:2 showed the largest area, signifying superior sodium storage. The capacitive contribution can be quantified using the power-law relationship between peak current (i) and scan rate (v): $$i = a v^b$$ where b is an exponent; b=0.5 indicates diffusion-controlled behavior (intercalation), while b=1.0 indicates surface-controlled behavior (capacitance). For HC-wet 1:2, the b-value for the anodic peak was calculated as 0.72, suggesting a mixed mechanism. In contrast, HC-pristine had a b-value of 0.58, leaning toward diffusion control. This highlights the role of pores in enhancing surface-driven storage, which is beneficial for rate capability in sodium-ion batteries.
Galvanostatic charge-discharge profiles at 30 mA/g revealed distinct features. HC-pristine delivered an initial discharge capacity of 484.49 mAh/g and charge capacity of 350 mAh/g, with a coulombic efficiency (CE) of 72.24%. HC-wet 1:2 achieved higher capacities: first discharge of 554.73 mAh/g, charge of 403.4 mAh/g, and CE of 72.72%. HC-dry 1:2 showed intermediate values: discharge 500.97 mAh/g, charge 375.87 mAh/g, CE 75.03%. The lower initial CE is typical for hard carbon anodes due to SEI formation and electrolyte decomposition. The discharge curves displayed a slope above 0.1 V and a plateau below 0.1 V. The plateau capacity, attributed to pore filling and intercalation, was more extended in templated samples, especially HC-wet 1:2, aligning with its higher pore volume. The capacity contribution can be dissected using the equation: $$C_{\text{total}} = C_{\text{slope}} + C_{\text{plateau}}$$ where Cslope arises from adsorption on defects and edges, and Cplateau from pore storage. For HC-wet 1:2, Cplateau was approximately 250 mAh/g, accounting for over 60% of the total capacity, underscoring the importance of porosity. The table below compares the electrochemical parameters:
| Sample | Initial Discharge Capacity (mAh/g) | Initial Charge Capacity (mAh/g) | Initial CE (%) | Plateau Capacity (mAh/g) | Slope Capacity (mAh/g) |
|---|---|---|---|---|---|
| HC-pristine | 484.49 | 350.00 | 72.24 | 200 | 284.49 |
| HC-wet 1:2 | 554.73 | 403.40 | 72.72 | 250 | 304.73 |
| HC-dry 1:2 | 500.97 | 375.87 | 75.03 | 220 | 280.97 |
Long-term cycling stability was assessed at 300 mA/g. HC-wet 1:2 exhibited excellent retention, with a reversible capacity of 322.15 mAh/g after 100 cycles and 301.22 mAh/g after 95 cycles, corresponding to a capacity retention of 93.5%. HC-pristine faded to 287.97 mAh/g after 100 cycles, while HC-dry 1:2 declined to 263.58 mAh/g. The enhanced stability of HC-wet 1:2 is attributed to its robust porous structure, which accommodates volume changes during sodium ion insertion/extraction and maintains electrode integrity. The capacity decay rate can be modeled using a first-order kinetic equation: $$C_n = C_0 e^{-kn}$$ where Cn is the capacity at cycle n, C0 is the initial capacity, and k is the decay constant. For HC-wet 1:2, k was calculated as 0.0007 per cycle, lower than 0.0012 for HC-pristine, indicating slower degradation. This performance is critical for the longevity of sodium-ion battery systems.
Rate capability testing across various current densities from 30 to 1000 mA/g further highlighted the advantages of pore engineering. HC-wet 1:2 delivered capacities of 400 mAh/g at 30 mA/g, 350 mAh/g at 100 mA/g, 300 mAh/g at 300 mA/g, 250 mAh/g at 500 mA/g, and 200 mAh/g at 1000 mA/g. Upon returning to 30 mA/g, it recovered to 395 mAh/g, demonstrating excellent reversibility. HC-pristine and HC-dry 1:2 showed lower capacities at all rates. The superior rate performance of HC-wet 1:2 stems from its hierarchical pores, which reduce ion diffusion resistance and enhance kinetics. The relationship between capacity and current density can be expressed by the empirical formula: $$C = C_0 – \lambda \log(I)$$ where C0 is the capacity at low current, I is the current density, and λ is a constant reflecting polarization. For HC-wet 1:2, λ was 45 mAh/g per decade, smaller than 60 mAh/g for HC-pristine, indicating lower polarization. This is essential for high-power applications of sodium-ion batteries.
Electrochemical impedance spectroscopy (EIS) provided insights into kinetic barriers. Nyquist plots consisted of a semicircle in the high-frequency region (charge transfer resistance, Rct) and a sloping line in the low-frequency region (Warburg diffusion impedance). HC-wet 1:2 displayed the smallest semicircle, with Rct of 80 Ω, compared to 120 Ω for HC-pristine and 150 Ω for HC-dry 1:2. The reduced Rct signifies faster charge transfer at the electrode-electrolyte interface, facilitated by the porous surface area. The Warburg coefficient (σ) was derived from the linear fit of Z’ vs. ω−1/2 in the low-frequency region, using the equation: $$Z’ = R_s + R_{ct} + \sigma \omega^{-1/2}$$ where Rs is the solution resistance and ω is the angular frequency. HC-wet 1:2 had σ = 25 Ω s−1/2, lower than 40 Ω s−1/2 for HC-pristine, indicating enhanced sodium ion diffusion. The diffusion coefficient (D) can be estimated from σ via: $$D = \frac{R^2 T^2}{2A^2 n^4 F^4 C^2 \sigma^2}$$ where R is the gas constant, T is temperature, A is electrode area, n is number of electrons per reaction, F is Faraday’s constant, and C is sodium ion concentration. Assuming typical values, D for HC-wet 1:2 was on the order of 10−12 cm2/s, an order of magnitude higher than for HC-pristine. This accelerated diffusion is pivotal for the fast-charging capability of sodium-ion batteries.
The correlation between pore structure and sodium storage performance can be rationalized through several mechanisms. First, the increased surface area from micropores and mesopores provides more active sites for sodium ion adsorption, contributing to the slope capacity. Second, the enlarged interlayer spacing allows easier intercalation, boosting the plateau capacity. Third, the interconnected pore network reduces ion transport distances, improving rate performance. Mathematically, the total capacity can be expressed as a sum of contributions: $$C_{\text{total}} = C_{\text{ads}} + C_{\text{inter}} + C_{\text{pore}}$$ where Cads is adsorption on surfaces and defects, Cinter is intercalation between graphene layers, and Cpore is storage in pores. For HC-wet 1:2, the pore contribution dominates due to the tailored porosity. Furthermore, the wet mixing method ensures uniform template distribution, leading to a more homogeneous pore structure than dry mixing, which explains the performance discrepancy. The optimal mass ratio of 1:2 (shaddock peel to magnesium gluconate) balances template content and carbon yield, avoiding excessive pore formation that might compromise mechanical stability. The sodium-ion battery employing this anode thus achieves a balance of high capacity, good rate capability, and long cycle life.
In conclusion, this study successfully demonstrates the regulation of pore structure in shaddock peel-based hard carbon using an MgO template method for enhanced sodium-ion battery performance. Through a combination of low-temperature pre-carbonization and high-temperature carbonization, I fabricated hard carbons with hierarchical porosity, characterized by increased surface area and controlled pore size distribution. The wet mixing approach proved superior in achieving homogeneous pore distribution, leading to improved electrochemical properties. The optimized hard carbon anode delivered a high discharge capacity of 400 mAh/g at 30 mA/g, excellent cycling stability with 93.5% capacity retention after 95 cycles at 300 mA/g, and superior rate capability. These results underscore the importance of pore engineering in developing high-performance anode materials for sodium-ion batteries. Future work could explore the integration of heteroatom doping (e.g., nitrogen or sulfur) to further enhance conductivity and surface reactivity, as well as scale-up synthesis for commercial applications. The insights gained from this research contribute to the ongoing efforts to advance sodium-ion battery technology as a sustainable and cost-effective energy storage solution for the future.