Sodium storage in fluorine-rich mesoporous carbon fabricated by low-temperature carbonization of polyvinylidene fluoride with a silica template

Abstract

A facile and manageable method has been developed to prepare fluorine-doped mesoporous carbon by carbonizing polyvinylidene fluoride, a fluorine-contained carbon precursor, with a silica template for the first time. And it is argued that the F contents depend directly on carbonization temperatures. A mesoporous carbon with a high F-content of 8.35 at% is obtained at a low carbonization temperature of 500 °C. The fluorine-rich mesoporous carbon (FMC) possesses an enlarged interlayer spacing of 0.402 nm, an enhanced electrical conductivity and a mesopore based structure, exhibiting a remarkable electrochemical energy storage performance evaluated as a NIB anode, such as a high stable reversible capacity of 385 mA h g⁻¹ at 100 mA g⁻¹, a satisfactory initial coulombic efficiency of 59% and a high rate capability of 116 mA h g⁻¹ at 5 A g⁻¹.

---

Introduction

With the rapidly increasing demand from the fields of low-cost energy storage systems and other large-scale applications, sodium ion batteries (NIBs) have recently attracted considerable attention as one of the most promising alternatives to lithium ion batteries (LIBs) due to the wide abundance and low cost of sodium resources.^1–7 It has been regarded as the major challenge to identify suitable cathode and anode materials.⁸ In the past few years, some oxide compounds, polyanion compounds, and other positive electrode compounds have been developed as cathode materials for NIBs, such as NaMO₂, NaMPO₄, NaMF₃ (M = Fe, Co, Ni, Mn, etc.).^9–16 The NIB anode materials can be classified into two categories: carbonaceous materials and non-carbonaceous materials. Among them, carbon materials have been widely used to fabricate LIB anodes for their abundance, low cost, environmental benignity, good electrical conductivity, excellent thermal and electrochemical stability.¹⁷ Since Li and Na share common properties in many aspects as alkali metals, it is seemly a wise means to look for good Na anode materials by starting at structures and chemistries that function well for Li intercalation. However, the ionic radii of Na (0.102 nm) is larger than that of Li (0.076 nm), which affects mass transport and storage during charge–discharge process and many of superior LIB anode materials are unsuitable for NIBs.¹⁸ Graphite is one of such anode materials, the small interlayer distance of which (∼0.34 nm) leads to its failing application as NIB anodes. As demonstrated by recent theoretical calculations, the minimum interlayer distance required for Na⁺ insertion in carbon materials is 0.37 nm.¹⁹

Porous hard carbons have been used widely as high-performance NIB anodes for the following reasons: (i) the high specific surface areas can offer large electrode/electrolyte interfaces for the charge-transfer reaction, and facilitate ion transport by providing shorter diffusion pathways; (ii) the low graphitization and high degree of disorder could provide more Na ion storage sites; (iii) the large interlayer distances are beneficial to Na⁺ insertion–extraction.^19–25 Especially, the mesoporous carbons (2 nm < pore size < 50 nm) which possess faster mass transport than that of microporous carbons (pore size < 2 nm) and higher specific surface areas compared with macroporous carbons (pore size > 50 nm), have been highly recommended for the fabrication of high-performance anodes for NIBs.^26–29

The porous carbon anodes for NIBs can be further improved by doping with heteroatoms such as B, N, P and S or co-doping with two or more kinds of such heteroatoms.^25,27–32 Among the heteroatoms, nitrogen has received more attention for its strong electronegativity and hybridization of nitrogen lone pair electrons with the π electrons in the carbon which can enhance the electronic conductivities of carbon anodes and increase binding sites for ion storage.^{25,29,32–36} In addition to nitrogen, fluorine is also used to improve electrochemical performance of carbon electrodes because of its higher electronegativity compared with nitrogen (χ = 3.98 of F versus 3.04 of N) and its semi-ionic bonding features.^37–44 Fluorinated graphite, fluorinated graphene and their derivatives have been intensely investigated and widely used as electrode materials for lithium ion batteries (LIBs) or supercapacitors.^37,39,40,43 Recently, the fluorine and nitrogen co-doped carbon microspheres were obtained and exhibited excellent capacitive performances employed for electrical double-layer capacitors.⁴⁵ The carbon-based materials were usually fluorinated by reacting with the fluorine-containing gases such as F₂, ClF₃, NF₃ or other fluorine-containing substances like hydrofluoric acid, ammonium fluoride (NH₄F) or ammonium fluoroborate (NH₄BF₄).^40–45 It may be wise to prepare fluorine-doped carbon by carbonizing directly fluorine-containing carbon precursors. And Wang et al.⁴⁶ fabricated fluorine-doped carbon particles by direct pyrolysis of the fluorine-containing biomass of lotus petioles, though its F-content was very low.

Polyvinylidene fluoride (PVDF) is a semi-crystalline polymer containing a crystalline phase and an amorphous and/or rubbery phase, with ultra-high fluoride content of about 60%. PVDF and its derivatives have been used to manufacture electrodes of energy-storage devices such as batteries and supercapacitors as binder, for their good chemical and thermal stabilities, excellent processability and high viscosity. PVDF were also used as fluorine-containing carbon sources to prepare F-doped carbon coating anode and cathode materials for LIBs.^47,48 Shi et al.⁴⁹ prepared fluorine-doped carbon anodes by carbonization of PVDF at 600–1000 °C for LIBs. The PVDF-derived carbons with the low fluorine content of 1.46–0.18% (at%) exhibit unsatisfactory electrochemical performances. In this report, we employ PVDF as a fluorine-containing carbon precursor to prepare ordered mesoporous fluorine-doped carbon by carbonization of PVDF with silica template at 500 °C for the first time (Scheme 1). The fabricated fluoride-doped mesoporous carbon (FMC) with pores centered at 3–4 nm and 10–20 nm, a surface area of 639.4 m² g⁻¹ and a F-content of 8.35 at% delivers a high initial reversible capacity up to 408 mA h g⁻¹ at a current density of 100 mA g⁻¹ with an initial coulombic efficiency of 59%, an excellent rate capability and cycling stability of 276 mA h g⁻¹ after 200 cycles at 1 A g⁻¹ as NIB anode. Even at an ultra-high current density of 5 A g⁻¹, a reversible capacity of 116 mA h g⁻¹ is still maintained.

Scheme 1 Schematic illustration of the synthesis procedure of FMC.

Results and discussion

It has been argued that the F contents depend directly on carbonization temperatures in the PVDF-derived carbons.⁴⁹ Fig. 1 shows the TG-DTG curves of PVDF, which indicate that the pyrolysis of PVDF starts from 400 °C and a sharp peak appears around 440 °C displaying the rapid weight loss and carbonization of the sample. The fluorine contents at different carbonization temperatures from 500 °C to 700 °C were investigated by XPS and XRF as listed in Table S1.† It can be found the F content of FMC gradually decreases with the increase of carbonization temperatures, which is consistent with the results reported in the literature.⁴⁹ The temperature of 500 °C is seemly the optimum condition under which PVDF is carbonized completely and the obtained FMC possesses the highest F content of 8.35 at%.

Fig. 1 TG-DTG curves of PVDF.

XPS was applied to investigate the nature of the fluorine species at the surface. XPS wide scan spectra of FMC with three strong peaks corresponding to C 1s, O 1s and F 1s are shown in Fig. 2a. The XPS spectrum of C 1s (Fig. 2b) can be divided into five peaks at 284.7, 285.7, 287.0, 288.5 and 290.7 eV. Among them, 284.7 eV corresponds to the sp² hybridized carbon (C–C), 285.7 eV is attributed to carbon atoms single bonded to oxygen in the form of epoxy, phenol or ether (C–O–C/C–OH), 287.0 eV belongs to the carbonyl carbon (C=O), 288.5 eV and 290.7 eV are associated with semi-ionic C–F bond and covalent C–F bond.⁴⁰ The relative atomic percentages and the binding energies are presented in Table 1. Although the precursor PVDF contains no oxygen, about 5.44 at% oxygen was detected in the FMC sample. A part of oxygen may be generated from the reaction of fluorine atoms with moisture of surrounding.⁵⁰ The O 1s spectrum contains two peaks at 531.9 and 533.5 eV, corresponding to the presence of different oxygen functionalities such as C=O and O–C–O/C–OH (Fig. 2c). These oxygen functionalities could result in the surface redox reaction between carbon–oxygen functional groups and Na⁺; i.e., –C=O + Na⁺ + e ↔ –C–O–Na, and provide extra ion storage sites.⁵¹ Moreover, the F 1s spectrum of the fluorinated sample has two peaks at 687.3 and 688.5 eV (Fig. 2d), which indicate nearly semi-ionic and covalent C–F bonds, respectively.⁴⁰ Fluorine possesses the strongest electronegativity among all the atoms which is believed to weaken the repulsion for Na ion insertion and extraction, lower the energy barrier for Na ion insertion and make favorable binding sites for Na⁺ storage in F-doped carbons evaluated as NIB anodes.⁴⁶ Moreover, the semi-ionic C–F bond could enhance the electrical conductivity of the carbon anodes which will benefit the improvement of their rate performance.⁵² The elemental mapping images demonstrate the uniform distribution of carbon, oxygen and fluorine in FMC, as shown in Fig. 2e.

Fig. 2 The total XPS spectrum (a) and the C 1s (b), the O 1s (c) and the F 1s (d) spectra and SEM image of a selected region and corresponding elemental maps of C, O and F (e) of FMC.

Table 1 Peak assignment of C 1s, O 1s and F 1s for FMC

Peak	Binding energy (eV)	Assignment	Fraction of species (%)
C 1s	284.7	C1: sp^2 C–C	72.5
	285.9	C2: C–O–C/C–OH	7.1
	287.0	C3: C=O	6.2
	288.3	C4: semi-ionic C–F	11.3
	290.8	C5: covalent C–F	2.9
O 1s	531.9	O1: C=O	41.3
	533.5	O2: O–C–O/C–OH	58.7
F 1s	687.3	F1: semi-ionic C–F	80.1
	688.5	F2: covalent C–F	19.8

---

EDS was applied to further analyze the content of elements in FMC shown in Fig. S1.† The EDS of FMC shows only the C, O and F elements without any other impurities and the relative contents of them are 86.15, 5.76 and 8.09 at%, respectively, which are coincident with the measure results of XPS and XRF. To observe directly the influence of F-doping on the surface property of PVDF-derived carbons, a simple contact angle measurement was carried out by using water as the droplets (Fig. 3). The average contact angle (CV) of FMC (130°, Fig. 3a) is larger than that of the fluoride-doped mesoporous carbon obtained at the higher calcination temperature of 600 °C (re-FMC) with low F-content of 2.17 at% (118°, Fig. 3b), indicating that the higher F-content brings about higher hydrophobic property. The bulk resistivity of FMC was measured by the four point probe method (Fig. S2†). The typical bulk resistivity of FMC was about 430 Ω cm, lower than that of the bulk resistivity of re-FMC (769 Ω cm). The better electronic conductivity of FMC could be contributed to its higher semi-ionic C–F content than that of re-FMC (Fig. S3 and Table S2†). Note that though the F-content of FMC is far higher than that of re-FMC, its resistivity is slightly higher than that of re-FMC, which may be because of the high-content oxygen in FMC. And the oxygen-containing functional groups will result in the decrease of resistivity and the bulk resistivity measured by the four point probe method is the result of synergistic effect of different heteroatom. Therefore, it is a challenge to prepare mesoporous carbons with high content of fluorine but low content of oxygen, which is to be paid more attention to during our following research.

Fig. 3 Photographs of water droplet on surface of FMC (a) and re-FMC (b).

As presented in the SEM image (Fig. 4a), the FMC consists mainly of three-dimensional (3D) short rods with a length of ∼500 nm and a width of ∼200 nm, and these rods arrange in a regular pattern along one certain direction similar to that of the ordered mesoporous silica template (Fig. S4a†). The microstructure could be further investigated by TEM and HR-TEM. Fig. 4b and c show most of short rods of FMC with linear arrays of mesopores and the bright contrast strips represent the pore wall images, indicating it possesses ordered mesopores with the pore diameter of about 4 nm measured directly from the partial enlargement of Fig. 4c (Fig. 4d). Some lager mesopores with pore diameter of ∼20 nm are observed on the edges of some rods as shown in Fig. 4e, the formation of which may be attributed to the not so good thermostability of PVDF derived carbon. During the carbonation process of PVDF/SBA-15 composite, splitting decomposition could take place along the occupied silica template in PVDF. After the removal of template, the mesopores in larger size would be obtained. The amorphous texture of FMC can be observed from the HR-TEM image (Fig. 4f), displaying its low degree of order.

Fig. 4 SEM image (a), TEM images (b), (c) and (e) and HR-TEM image (f) of FMC, (d) is the partial enlargement of (c).

Nitrogen adsorption–desorption isotherm was employed to further characterize the porous structure of the FMC sample (Fig. 5a). The FMC exhibits a type IV isotherm with a distinct hysteresis loop at higher relative pressures (P/P₀ > 0.5), indicating the presence of well ordered mesopores.⁵³ The specific surface area of FMC, calculated by multipoint BET, was 639.4 m² g⁻¹ and the total pore volume was 1.356 cm³ g⁻¹, mainly contributed by the mesoporous pores (∼89%, by t-plot method). The pore size distribution (PSD) was calculated using the non-local density function theory (NLDFT) model, as shown in Fig. 5b. There is a weak peak within the micropore range, a sharp peak centered at 3.9 nm and a weak and wide hump located at 10–20 nm. This indicates that the FMC carbon consists of mainly two kinds of mesopores and a small amount of micropores. The formation mechanism of micropores is based on the release of HF from the carbon chain in PVDF during carbonization under insert atmosphere.⁴⁹ The result of mesoporous structure analysis is consistent with the TEM measurement. This special porous structure is important for the improved electrochemical performance of FMC anode for NIBs because the mesopores can serve as Na⁺ reservoirs and pathways for the electrolyte, while the micropores are suggested as charge accommodation which will be essential for high energy storage.^27,29,54 The PVDF-derived carbon obtained without template at 500 °C delivers a specific surface area of 113.2 m² g⁻¹ and a non-mesoporous smooth surface as shown in Fig. S4b,† indicating the use of mesoporous silica template is critical to achieve the high specific surface area and porous structure.

Fig. 5 Nitrogen adsorption–desorption isotherm (a), NLDFT pore-size distribution (b), XRD spectrum (c) and Raman spectra (d) of FMC. The inset of (c) shows a schematic depicting of R.

The crystallite structure of FMC was characterized by XRD and Raman spectroscopy. As shown in Fig. 5c, the XRD pattern demonstrates two broad diffraction peaks at 21.2° and 42.8° that are indexed as (002) and (100) of the pseudographitic domains. An empirical parameter (R), defined as the ratio of height of the (002) Bragg peak to the surrounding background, was used to signify the graphitization degree of the material.⁵⁵ The inset of Fig. 5c shows a schematic depicting how R was calculated. The R value of FMC is 1.96 displaying the not so high degree of graphitization in accord with the HR-TEM measurement. And the relatively amorphous structure could accommodate more Na ions for NIBs applications.⁵⁶ The interlayer spacing (d₀₀₂) value of the FMC was calculated by Bragg equation based on the (002) peak to be 0.402 nm. The interlayer distance of FMC is much larger than the minimum interlayer distance for Na ion insertion (0.37 nm) as predicted by theoretical calculations,³⁷ which is beneficial for Na⁺ intercalation and deintercalation. The enlarged interlayer distance can be ascribed to the F-doping.^43,46 The Raman spectroscopy of the FMC (Fig. 5d) shows a wide D-band at ∼1345 cm⁻¹ and a G-band at 1591 cm⁻¹, respectively. It is well known that the D-band can be ascribed to edges, other defects, and disordered carbon, while the G-band is the vibration of sp²-bonded carbon atoms in a 2D hexagonal lattice in carbon materials.⁵⁷ And the disorder degree of the carbon can be estimated using the I_D/I_G ratio. The I_D/I_G value of FMC is ∼0.81, much higher than that of commercial graphite (0.37),⁵⁸ indicating it possesses a high degree of disorder and some defects caused by the heteroatom of F or/and O which could provide extra ion-storage sites.

It is worthy to note that FMC delivers a higher value of R and a lower I_D/I_G value displaying a little higher degree of order, compared with the biomass-derived porous carbons reported in other works of our research group,^59,60 which can be attributed to the semi-crystalline structure of PVDF containing a crystalline phase and an amorphous phase (Fig. S5†). During carbonization, the crystalline phase is beneficial to improving the ordering degree of FMC.

According to the points above, the FMC with high content of fluorine, mesopore-based and moderately disordered structure is expected to be a promising anode material for sodium ion battery. The electrochemical performances of FMC as NIB anode were investigated in half-cell configurations countered with metallic sodium. As shown in Fig. 6a, the FMC anode exhibits typical CV curves of carbon anode materials with a pronounced and broad cathodic peak at 0.01–1 V during the first cycle, which can be ascribed to the irreversible consumption of charge caused by the formation of the solid electrolyte interphase (SEI) layer and the loss of Na⁺ in the irreversible storage sites. Although the intensity of this cathodic peak in the first cycle is stronger than that in the following two cycles, the second and third CV curves almost overlap, indicating that Na⁺ insertion–extraction in FMC electrode is very stable.

Fig. 6 Cyclic voltammograms of FMC at a scan rate of 0.1 mV s⁻¹ (a), charge–discharge curves of FMC at 0.1 A g⁻¹ (b), capacity over cycling at discharge–charge current densities of MC and FMC (c) and typical Nyquist plots of MC and FMC (d).

Fig. 6b shows the charge–discharge profiles of the FMC electrode between 0.01 and 2.8 V. All the discharge curves consist of a slope in a higher potential range attributed to Na⁺ insertion into the enlarged carbon layers with an interlayer spacing of 0.402 nm and an approximative plateau in the narrow low potential range which can be ascribed to Na⁺/Na filling in the small mesopores with a pore diameter of about 4 nm and some micropores, which is reconciled well with the mechanism model of Na⁺ storage in the amorphous carbon.^18,61,62 The first charge–discharge cycle reveals the initial discharge and charge capacities to be 692 and 408 mA h g⁻¹ at a current density of 0.1 A g⁻¹, showing an initial coulombic efficiency of 59%. The irreversible capacity loss mainly originates from the formation of SEI layer and Na⁺ insertion in the irreversible storage sites, coinciding with the results from the CV observation. It is well known that the porous carbon materials as anodes for NIBs tend to exhibit very low initial coulombic efficiency due to the larger ionic radii of Na versus Li, needing larger interlayer spacing for the reversible insertion–extraction of Na⁺. In addition, this larger ionic radii of Na is easy to cause its capture in the carbon anodes. However, FMC delivers a higher initial efficiency, compared with the porous carbon anodes reported in the literatures.^{20,21,23,25,27,29}

The improved initial efficiency of FMC can be ascribed to the F-doping which can reduce the electrochemical decomposition of ethylene carbonate (EC) and the loss of Na⁺ during the formation of SEI film on the one hand, and enlarge the interlayer distances of carbon materials which will make for the intercalation and deintercalation of Na ions and avoid the capture of lots of Na⁺ in carbon layers on the other hand.^43,46 After 10 cycles, the reversible capacity of the FMC stabilizes at 385 mA h g⁻¹ and the stable reversible is satisfactory.

In order to compare with FMC, the mesoporous carbon without fluorine (MC) has been obtained by carbonizing glucose with SBA-15. Fig. 6c shows the capacities of MC and FMC at various discharge–charge current densities during cycling. The FMC electrode delivered stable reversible capacities of 385, 341, 318, 286, 196 and 116 mA h g⁻¹ at current rates of 0.1, 0.3, 0.5, 1, 3 and 5 A g⁻¹, respectively. When the charge–discharge current density turns back to 0.1 A g⁻¹, the reversible capacity of FMC can be recovered to 375 mA h g⁻¹. Obviously, the reversible capacities of FMC at all the current densities are higher than those of MC. When the current densities increase from 0.1 A g ⁻¹ to 5 A g ⁻¹, the stable reversible capacities of MC decrease from 236 mA h g⁻¹ to 49 mA h g⁻¹ with a higher capacity loss ratio than that of FMC, indicating FMC possesses superior rate capability to MC. The higher capacity and better rate capability of FMC can be ascribed to its high F-doping. Fig. 6d shows the impedance spectra of the MC and FMC, with an equivalent circuit model in the inset. The intercept with the real impedance (Z′) axis in the high frequency region can be related to the ohmic resistance (R_e), corresponding to the intercept with the X-axis in high frequency, including the intrinsic resistance of the active material and the electrolyte, as well as the contact resistance at the interface between the active material and the current collector.²⁹ The semicircle corresponds to surface film resistance and charge transfer at the interface between the electrode and electrolyte (R_sf+ct), while the inclined line accounts for characteristic Warburg resistance (Z_w) related to the mass transfer process within the pores of the carbon.²⁸ Moreover, a constant phase element (CPE) is introduced in order to precisely simulate the capacitance contributed by the surface of the active material. Comparing the Nyquist plots of MC and FMC, it is found the FMC electrode exhibits a lower R_sf+ct value (211 Ω) than that of MC (304 Ω), which implies that FMC possesses lower SEI resistance, higher electrical conductivity and faster charge transfer reaction for Na⁺ insertion/extraction than MC.

Especially, the reversible capacity of FMC (116 mA h g⁻¹) at 5 A g⁻¹ is extremely high among the carbon anodes reported recently.^18–23 Although the slightly better rate performances of some carbon materials as NIB anodes have been published, these carbon materials were prepared either by sophisticated processes or/and needing expensive reagents.^17,24,30,32 In this paper, we employed a low-cost fluorine contained precursor of PVDF to fabricate the fluorine-doping mesoporous carbon with prominent rate performances by a simple low-temperature carbonization method, which is meaningful in mass production and application.

The main reasons for the excellent rate capability can be concluded as follows. Firstly, F-doping can enlarge the interlayer distance of carbon anode which is beneficial to fast insertion/extraction of Na⁺ as depicted in Fig. 7. Secondly, the mesopore-based structure offer not only large amounts of pores serving as Na-ion reservoirs, but also lots of pathways for the electrolyte diffusion which can accelerate the kinetic process of ion diffusion in the electrodes and improve the rate performances.⁶³ Lastly, the slightly higher degree of order compared with other hard carbons also make for the improvement of electrical conductivity and rate performance.

Fig. 7 Schematic representation Na ion storage in FMC.

The highly electronegative fluorine and strong single covalent bond of C–F can also enhance charge polarization and reduce the overpotentials, which will improve electrochemical activity of energy-related reactions and electrode cycling stability.⁶⁴ The cycling performance and corresponding coulombic efficiency of FMC for 200 charge/discharge cycles at a current density of 1 A g ⁻¹ were shown in Fig. 8. The reversible capacity is still retained at 276 mA h g⁻¹ after 200 cycles, and the coulombic efficiency exceeds 97% after the first cycle, representing good cyclic stability and reversibility of the FMC anode.

Fig. 8 Cycling performance and corresponding coulombic efficiency of FMC at a current density of 1 A g⁻¹.

Conclusions

The fluoride-doped mesoporous carbon (FMC) with a high F-content and mesopore based structure has been fabricated by low-temperature carbonization of PVDF with silica template, followed by the template removal. The FMC carbon as NIB anode exhibits a high reversible capacity of 385 mA h g⁻¹ at 100 mA g⁻¹, a satisfactory initial coulombic efficiency of 59%, a high rate performance of 116 mA h g⁻¹ at 5 A g⁻¹ and an excellent cycling stability with a retained capacity of 276 mA h g⁻¹ after 200 cycles at 1 A g⁻¹. The excellent electrochemical performances can be attributed to F-doping and special porous structure consisting of large amounts of mesopores and some micropores. F-Doping can enlarge the interlayer spacing of carbon materials which will make for the intercalation and deintercalation of Na⁺, enhance the electrical conductivity of the carbons which will benefit the improvement of rate performance and reduce the decomposition of EC to improve the initial coulombic efficiency. The mesopores in FMC can serve as Na⁺ reservoirs and pathways for the electrolyte which will improve further the capacity and rate capability, while the micropores are suggested as charge accommodation and essential for high energy storage in NIB application.

Experimental

Preparation of fluoride-doped mesoporous carbon (FMC)

As shown in Scheme 1, 1 g PVDF was first dissolved in 20 mL N-methyl-2-pyrrolidone (NMP) to form a transparent solution, and then 1 g silica template (SBA-15) with pore size 4–6 nm (purchased from Nasenmei nano materials, Hangzhou, China) was added into the solution and mixed sufficiently. The mixture was under stirring for 12 h and dried at 120 °C, followed by carbonized in an argon-flowing tube furnace at 500 °C for 2 h with a heating rate of 5 °C min⁻¹. 2 mol L⁻¹ NaOH aqueous solution was used to absorb the gas of HF released during the carbonization process. Sintered products were treated in 10% HF solution for 6 h to remove the silica template, washed with double-distilled water to neutral, and dried at 90 °C under vacuum for 12 h. The as-prepared sample was abbreviated to FMC. For comparison, the fluoride-doped mesoporous carbon obtained at the higher calcination temperature of 600 °C was designated as re-FMC.

General characterization

The TG-DTG analysis was performed using a SDT Q600 thermal analyzer (New Castle, USA). X-ray photoelectron spectra (XPS) were recorded by using a Kratos XSAM 800 spectrometer (Manchester, UK). X-ray fluorescence (XRF) spectra were acquired by an Axios XRF spectrometer, PANalytical B. V. (Almelo, The Netherlands). Field emission scanning electron microscopy (SEM) images and energy dispersive spectrum (EDS) were obtained with a Hitachi S4800 scanning electron microscope (Tokyo, Japan). Transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HR-TEM) images were observed on a FEI Tecnai G2 20 TEM (Hillsboro, OR, USA) operating at 200 kV. A Micromeritics TriStar II 3020 automatic analyzer (Norcross, GA, USA) was employed to carry out nitrogen sorption analysis. The sample was degassed at 120 °C for 24 h before absorption measurements. Specific surface areas were estimated according to the BET model and pore size distributions (PSDs) were calculated by applying a density function theory (DFT) model. A TD-3500 X-ray powder diffractometer (Danton, China) was used to perform X-ray diffraction (XRD) measurements. Raman spectra were detected by using a Confocal LabRAM HR800 spectrometer, HORIBA Jobin Yvon (Paris, France). The average contact angle (CA) values were obtained by measuring the same film at three different positions using the sessile drop method of a OCAH200 contact angle goniometer, Dataphysics (Stuttgart, Germany). The bulk resistivity was determined using a ST-2722 semiconductor resistivity of the power tester (Suzhou, China).

Electrochemical measurements

The electrochemical behavior of the FMC anode was evaluated using 2032-type coin cells. To prepare working electrodes, the slurry of active material, acetylene black (conductive additive) and PVDF (binder) at a weight ratio of 80 : 10 : 10 in N-methyl-2-pyrrolidone (NMP) was coated onto a copper foil current collector and then dried at 110 °C overnight in vacuum oven. The obtained working electrode, with a loading of around 1.2 mg cm⁻², were assembled in an argon-filled glove box using sodium metal for the counter electrodes. The electrolyte was 1 M NaClO₄ in 1 : 1 by volume ethylene carbonate (EC) and diethyl carbonate (DEC). Cyclic voltammetry (CV) measurements were performed at a scan rate of 0.1 mV s⁻¹ on an Autolab PGSTAT 302 electrochemical workstation (Utrecht, the Netherlands) in the voltage range of 0.01–2.8 V (vs. Na⁺/Na). Electrochemical impedance spectroscopy (EIS) was carried out in the frequency range 0.01–100 kHz at a charged stage with an applied amplitude of 5 mV on the same electrochemical workstation. The galvanostatic charge–discharge tests were carried out between 0.01 and 2.8 V at room temperature on a Neware CT-3008W battery cycler (Guangdong, China).

Acknowledgements

The present work was supported by National Natural Science Foundation of China (No. 21506131) and Science and Technology Support Program of Sichuan Province (No. 2014GZ0095 and 2015RZ0057).

Notes and references

1. H. Pan, Y.-S. Hu and L. Chen, Energy Environ. Sci., 2013, 6, 2338–2360.
2. V. Palomares, P. Serras, I. Villaluenga, K. B. Hueso, J. Carretero-González and T. Rojo, Energy Environ. Sci., 2012, 5, 5884–5901.
3. M. D. Slater, D. Kim, E. Lee and C. S. Johnson, Adv. Funct. Mater., 2013, 23, 947–958.
4. N. Yabuuchi, K. Kubota, M. Dahbi and S. Komaba, Chem. Rev., 2014, 114, 11636–11682.
5. D. Kundu, E. Talaie, V. Duffort and L. F. Nazar, Angew. Chem., Int. Ed., 2015, 54, 3431–3448.
6. Z.-W. Zhou, Y.-T. Liu, X.-M. Xie and X.-Y. Ye, Chem. Commun., 2016, 52, 8401–8404.
7. Z.-W. Zhou, Y.-T. Liu, X.-M. Xie and X.-Y. Ye, ACS Appl. Mater. Interfaces, 2016, 8, 7092–7100.
8. S. W. Kim, D. H. Seo, X. Ma, G. Ceder and K. Kang, Adv. Energy Mater., 2012, 2, 710–721.
9. R. Berthelot, D. Carlier and C. Delmas, Nat. Mater., 2011, 10, 74–80.
10. N. Yabuuchi, M. Kajiyama, J. Iwatate, H. Nishikawa, S. Hitomi, R. Okuyama, R. Usui, Y. Yamada and S. Komaba, Nat. Mater., 2012, 11, 512–517.
11. M. Sathiya, K. Hemalatha, K. Ramesha, J.-M. Tarascon and A. Prakash, Chem. Mater., 2012, 24, 1846–1853.
12. K. T. Lee, T. Ramesh, F. Nan, G. Botton and L. F. Nazar, Chem. Mater., 2011, 23, 3593–3600.
13. C. Zhu, K. Song, P. A. van Aken, J. Maier and Y. Yu, Nano Lett., 2014, 14, 2175–2180.
14. B. Ellis, W. Makahnouk, Y. Makimura, K. Toghill and L. Nazar, Nat. Mater., 2007, 6, 749–753.
15. Y. Yamada, T. Doi, I. Tanaka, S. Okada and J.-I. Yamaki, J. Power Sources, 2011, 196, 4837–4841.
16. J. Song, L. Wang, Y. Lu, J. Liu, B. Guo, P. Xiao, J.-J. Lee, X.-Q. Yang, G. Henkelman and J. B. Goodenough, J. Am. Chem. Soc., 2015, 137, 2658–2664.
17. J. Xu, M. Wang, N. P. Wickramaratne, M. Jaroniec, S. Dou and L. Dai, Adv. Mater., 2015, 27, 2042–2048.
18. Y. Wen, K. He, Y. Zhu, F. Han, Y. Xu, I. Matsuda, Y. Ishii, J. Cumings and C. Wang, Nat. Commun., 2014, 5, 4033–4042.
19. Y. Cao, L. Xiao, M. L. Sushko, W. Wang, B. Schwenzer, J. Xiao, Z. Nie, L. V. Saraf, Z. Yang and J. Liu, Nano Lett., 2012, 12, 3783–3787.
20. Y. Yan, Y. X. Yin, Y. G. Guo and L. J. Wan, Adv. Energy Mater., 2014, 4, 1301584–1301588.
21. S. Wenzel, T. Hara, J. Janek and P. Adelhelm, Energy Environ. Sci., 2011, 4, 3342–3345.
22. S. R. Prabakar, J. Jeong and M. Pyo, Electrochim. Acta, 2015, 161, 23–31.
23. K.-L. Hong, L. Qie, R. Zeng, Z.-Q. Yi, W. Zhang, D. Wang, W. Yin, C. Wu, Q.-J. Fan and W.-X. Zhang, J. Mater. Chem. A, 2014, 2, 12733–12738.
24. H. Hou, C. E. Banks, M. Jing, Y. Zhang and X. Ji, Adv. Mater., 2015, 27, 7861–7866.
25. D. Li, H. Chen, G. Liu, M. Wei, L.-X. Ding, S. Wang and H. Wang, Carbon, 2015, 94, 888–894.
26. C. Liang, Z. Li and S. Dai, Angew. Chem., Int. Ed., 2008, 47, 3696–3717.
27. S. Zhang, F. Yao, L. Yang, F. Zhang and S. Xu, Carbon, 2015, 93, 143–150.
28. W. Li, Z. Yang, M. Li, Y. Jiang, X. Wei, X. Zhong, L. Gu and Y. Yu, Nano Lett., 2016, 16, 1546–1553.
29. H. Liu, M. Jia, N. Sun, B. Cao, R. Chen, Q. Zhu, F. Wu, N. Qiao and B. Xu, ACS Appl. Mater. Interfaces, 2015, 7, 27124–27130.
30. W. Li, M. Zhou, H. Li, K. Wang, S. Cheng and K. Jiang, Energy Environ. Sci., 2015, 8, 2916–2921.
31. J. P. Paraknowitsch and A. Thomas, Energy Environ. Sci., 2013, 6, 2839–2855.
32. D. Xu, C. Chen, J. Xie, B. Zhang, L. Miao, J. Cai, Y. Huang and L. Zhang, Adv. Energy Mater., 2016, 6, 1501929–1501935.
33. C. Yang, W. Li, Z. Yang, L. Gu and Y. Yu, Nano Energy, 2015, 18, 12–19.
34. F. Yang, Z. Zhang, K. Du, X. Zhao, W. Chen, Y. Lai and J. Li, Carbon, 2015, 91, 88–95.
35. Z. Guan, H. Liu, B. Xu, X. Hao, Z. Wang and L. Chen, J. Mater. Chem. A, 2015, 3, 7849–7854.
36. Z. Wang, Y. Li and X.-J. Lv, RSC Adv., 2014, 4, 62673–62677.
37. T. Nakajima, V. Gupta, Y. Ohzawa, M. Koh, R. N. Singh, A. Tressaud and E. Durand, J. Power Sources, 2002, 104, 108–114.
38. Y.-S. Lee, J. Fluorine Chem., 2007, 128, 392–403.
39. R. R. Nair, W. Ren, R. Jalil, I. Riaz, V. G. Kravets, L. Britnell, P. Blake, F. Schedin, A. S. Mayorov and S. Yuan, small, 2010, 6, 2877–2884.
40. M.-J. Jung, E. Jeong, S. Kim, S. I. Lee, J.-S. Yoo and Y.-S. Lee, J. Fluorine Chem., 2011, 132, 1127–1133.
41. T. Nakajima, J. Fluorine Chem., 2013, 149, 104–111.
42. M.-H. Kim, J.-H. Yang, Y.-M. Kang, S.-M. Park, J. T. Han, K.-B. Kim and K. C. Roh, Colloids Surf., A, 2014, 443, 535–539.
43. S. Huang, Y. Li, Y. Feng, H. An, P. Long, C. Qin and W. Feng, J. Mater. Chem. A, 2015, 3, 23095–23105.
44. T. Nakajima, J. Fluorine Chem., 2007, 128, 277–284.
45. J. Zhou, J. Lian, L. Hou, J. Zhang, H. Gou, M. Xia, Y. Zhao, T. A. Strobel, L. Tao and F. Gao, Nat. Commun., 2015, 6, 8503–8510.
46. P. Wang, B. Qiao, Y. Du, Y. Li, X. Zhou, Z. Dai and J. Bao, J. Phys. Chem. C, 2015, 119, 21336–21344.
47. H. Geng, Q. Zhou, Y. Pan, H. Gu and J. Zheng, Nanoscale, 2014, 6, 3889–3894.
48. L. Li, C. Fan, X. Zhang, T. Zeng, W. Zhang and S. Han, New J. Chem., 2015, 39, 2627–2632.
49. L. Shi, C. L. Miao, G. R. Chen, B. Xu and S. Chen, Adv. Mater. Res., 2012, 463–464, 730–733.
50. T. Nakajima and K. Yanagida, Tanso, 1996, 174, 195–200.
51. Y. Shao, J. Xiao, W. Wang, M. Engelhard, X. Chen, Z. Nie, M. Gu, L. V. Saraf, G. Exarhos and J.-G. Zhang, Nano Lett., 2013, 13, 3909–3914.
52. J. H. Lee, G. K. W. Koon, D. W. Shin, V. Fedorov, J. Y. Choi, J. B. Yoo and B. Özyilmaz, Adv. Funct. Mater., 2013, 23, 3329–3334.
53. Y. Qu, Z. Zhang, X. Zhang, G. Ren, Y. Lai, Y. Liu and J. Li, Carbon, 2015, 84, 399–408.
54. D. W. Wang, F. Li, M. Liu, G. Q. Lu and H. M. Cheng, Angew. Chem., 2008, 120, 379–382.
55. Y. Liu, J. Xue, T. Zheng and J. Dahn, Carbon, 1996, 34, 193–200.
56. J. Ding, H. Wang, Z. Li, K. Cui, D. Karpuzov, X. Tan, A. Kohandehghan and D. Mitlin, Energy Environ. Sci., 2015, 8, 941–955.
57. A. C. Ferrari and J. Robertson, Phys. Rev. B: Condens. Matter Mater. Phys., 2000, 61, 14095–140107.
58. B. Fang, M.-S. Kim, J. H. Kim, S. Lim and J.-S. Yu, J. Mater. Chem., 2010, 20, 10253–10259.
59. L. Chen, Y. Zhang, C. Lin, W. Yang, Y. Meng, Y. Guo, M. Li and D. Xiao, J. Mater. Chem. A, 2014, 2, 9684–9690.
60. J. Ou, Y. Zhang, L. Chen, Q. Zhao, Y. Meng, Y. Guo and D. Xiao, J. Mater. Chem. A, 2015, 3, 6534–6541.
61. D. Stevens and J. Dahn, J. Electrochem. Soc., 2001, 148, A803–A811.
62. D. Stevens and J. Dahn, J. Electrochem. Soc., 2000, 147, 1271–1273.
63. Z. Li, Z. Xu, X. Tan, H. Wang, C. M. Holt, T. Stephenson, B. C. Olsen and D. Mitlin, Energy Environ. Sci., 2013, 6, 871–878.
64. I. Y. Jeon, M. J. Ju, J. Xu, H. J. Choi, J. M. Seo, M. J. Kim, I. T. Choi, H. M. Kim, J. C. Kim and J. J. Lee, Adv. Funct. Mater., 2015, 25, 1170–1179.

---

Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra24386j

---