Three-dimensional sandwich-type graphene@microporous carbon architecture for lithium–sulfur batteries

Abstract

The commercial applications of lithium–sulfur batteries are hindered by several issues including the poor electronic/ionic conductivity of sulfur and discharge products, the dissolution of lithium polysulfides in organic electrolytes, and the volume change during charge/discharge processes. In this study, a three-dimensional (3-D) sandwich-type graphene@microporous carbon (G@MC) architecture with large pore volume (2.65 cm³ g⁻¹) and ultrahigh surface area (3374 m² g⁻¹) was designed to encapsulate sulfur and polysulfides in the hierarchical microporous structure. The G@MC materials with a lot of sp² hybrid carbon atoms can provide 3-D electron transfer pathways for sulfur and discharge products. Furthermore, the G@MC materials with the novel hierarchical structure can absorb a lot of polysulfides and restrain the polysulfide diffusion, and provide adequate nanospace for sulfur expansion ensuring the structural integrity during the cycling. Thus, the optimized G@MC–S nanocomposite with high sulfur loading (75.4 wt%) retains a discharge capacity of 541.3 mA h g⁻¹ after 500 cycles at 0.5C. This design strategy is simple and broadly applicable, providing new opportunities for materials design that can be extended to various electrode materials.

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Introduction

The lithium–sulfur battery with high theoretical specific capacity (1675 mA h g⁻¹) and high energy density (2600 W h kg⁻¹) is considered as one of the most promising candidates for advanced energy storage devices.^1–3 Furthermore, sulfur also possesses other advantages such as natural abundance, low cost, and non-toxicity. However, the commercial application of lithium–sulfur batteries has not been very successful because of the highly electrically and ionically insulating of sulfur and discharge products (Li₂S and Li₂S₂), the dissolution of lithium polysulfides in organic electrolytes, and the volume expansion of sulfur during discharge.⁴ The dissolution of polysulfides in organic electrolyte can result in the troublesome shuttling loss during cycling. It decreases active mass utilization and markedly reduces coulombic efficiency.⁵ Various strategies have been developed to address these issues, including the fabrication of porous carbon–sulfur composites,^6–16 preparation of polymer–sulfur composites,^17–19 optimization of organic electrolyte,^20,21 and use of coating layers^22–26 or carbon interlayers.^27–29 Recently, micropores were proven to be the most effective pore structure in confining the polysulfides diffusion.^12,30,31 However, the low pore volume leads to the low sulfur loading in the microporous carbon. Furthermore, the microporous carbon constituted by amorphous carbon results in the poor rate performance.²⁴ Thus, how to simultaneously improve the pore volume and the conductivity of microporous carbon is crucial to the development of the microporous carbon–sulfur nanocomposites.

Graphene constituted by sp² hybrid carbon has been considered as one of the most promising conductive matrixes for lithium–sulfur batteries because of its unique two-dimensional structure, excellent conductivity, high surface area, chemical stability, and flexibility.³² The graphene-based sulfur-containing cathode materials have excellent rate performance. However, the graphene and graphene sheets cannot restrain the diffusion of polysulfides effectively.^22,33 Thus, many graphene derivatives such as reduced graphene oxides,^34–36 N-doped graphene sheets,³⁷ chemical functionalized graphene,³⁸ and coating layer²² were used to confine the polysulfides diffusion. Most recently, graphene and porous carbon hybrid materials were proven as promising matrix for sulfur cathodes.^39,40 It is difficult to increase the percentage of microporous structure and improve the sulfur containing in these graphene based hybrid material simultaneously.

Herein, we design a novel three-dimensional (3-D) sandwich-type graphene@microporous carbon (G@MC) structure for lithium sulfur battery (Scheme 1). By combining the novel hierarchal microporous structure and graphene sheets, the G@MC materials achieve large pore volume (2.65 cm³ g⁻¹), high percentage of micropore, and high conductivity simultaneously. Furthermore, the optimized G@MC material can absorb a lot of polysulfides. Thus, the optimized G@MC–sulfur (G@MC–S) nanocomposite can deliver a capacity of 541.3 mA h g⁻¹ at a high sulfur loading (75.4 wt%) after 500 cycles at 0.5C (1C = 1675 mA h g⁻¹).

Scheme 1 Schematic illustration of the G@MC–S nanocomposite for lithium–sulfur battery.

Experiments

Material preparation

Graphene oxide (GO) was synthesized by oxidation of graphite using modified hummers method.⁴¹ 15 g of sucrose was added to 200 mL of GO aqueous solution (0.75 mg mL⁻¹) under sonication, which was placed in round bottom flask. Then, 100 mL sulfuric acid was added to the solution and refluxed at 120 °C for 10 h. The resulting black suspension was then filtered several times and dried at 100 °C for 12 h. The obtained material was infiltrated with KOH with mass ratio of 1 : 4 via soaking in a solution followed by stirring for 12 h. The impregnated slurry was dried at 110 °C for 12 h and then heated at 900 °C for 2 h in a tube furnace under Ar atmosphere with a heating rate of 5 °C min⁻¹. Subsequently, the obtained product was neutralized with 1 M HCl solution and washed with distilled water for several times. Finally, the carbon was dried at 100 °C for 20 h. The obtained product was nominated as G@MC4 (4 denotes the KOH to carbon mass ratio in the active process). The G@MC2, G@MC3, G@MC5, and G@MC6 samples were prepared as well according to similar route mentioned above.

The 150 mg of G@MC4 material was mixed with sublimed sulfur (350 mg) and was ground for 30 min in an agate mortar. Afterwards, the mixture was sealed in an evacuated quartz tube and heated at 155 °C for 6 h. Then the temperature was increased to 400 °C and kept at this temperature for 10 h and cooled down to room temperature. The obtain product was nominated as G@MC4–S nanocomposite. Additionally, using different carbon precursor in the route mentioned above, the G@MC2–S, G@MC3–S, G@MC5–S, G@MC5–S–H (using 500 mg sublimed sulfur), and G@MC6–S samples were obtained.

Material characterization

X-ray powder diffraction (XRD) patterns were obtained on a Rigaku D/max-2500 (Cu Kα radiation, λ = 0.15405 nm) operating at 5° min⁻¹. The microstructure of the samples was examined with a JEOL 6710F field-emission scanning electron microscope (FE-SEM) and a Tecnai G2 F20 U-TWIN field-emission transmission electron microscope (FE-TEM). The N₂ adsorption–desorption analysis was performed using an Autosorb-1 analyzer from Quantachrome Instruments. Raman spectra were obtained using a DXR from Thermo Scientific with a laser wavelength of 532 nm. Thermal analysis was measured on a TG/DTA 6300 instrument, in which the sample was heated in alumina crucible under N₂ flow to 500 °C at a heating rate of 10 °C min⁻¹. X-ray photoelectron spectroscopy (XPS) was performed on the Thermo Scientific ESCALab 250Xi using 200 W monochromated Al Kα radiation.

Electrochemical test

The cathode slurry was prepared by mixing 75 wt% G@MC–S nanocomposite, 15 wt% Ketjen black, and 10 wt% of polyvinylidene difluoride (PVDF, Alfa Aesar) dissolved in N-methyl-2-pyrrolidone (NMP, Aldrich). The sulfur cathodes were produced by coating the slurry on aluminum foil and drying at 60 °C for 12 h. The cell tests were evaluated using coin cells cycled at room temperature between 1.8 V and 2.7 V, which were fabricated in an argon-filled glove box using lithium metal as the counter electrode and a microporous polyethylene separator. The electrolyte was 1 M bis-(trifluoromethane)sulfonamide lithium (LiTFSI) in a mixed solvent of 1,2-dimethoxyethane and 1,3-dioxolane (1 : 1, v/v) with 0.1 M LiNO₃ additives. On average, the electrode thickness is ∼25 μm, the electrode diameter is 10 mm, and the areal sulfur mass loading is ∼1.5 mg cm⁻². The electrolyte for each cell is 20 μL per mg-sulfur, and the thickness of the metallic lithium anode is about 110 μm. The performance of the cells was tested using a LAND electrochemical testing system. The specific capacity was calculated on the mass of elemental sulfur.

Results and discussion

GO was used as the starting materials, then the carbon layers were generated on the graphene oxides layers during the first carbonization process with the assistance of sulfuric acid. The activation process yielded a continuous 3D network of micropores, and the carbon layers were assembled into sandwich-type multilayer architecture. The sulfur with the lowest viscosity at 155 °C could easily diffuse into the microporous structure of G@MC materials. Furthermore, sulfur on the external surface of the G@MC materials was sublimed at 400 °C, meanwhile the chemical interaction between sulfur and carbon was enhanced.⁴²

Fig. 1 shows the SEM and TEM images of typical G@MC and G@MC–S materials. The G@MC5 material is stacked by wrinkled multilayer nanosheets, and the particle size of the G@MC5 material can be up to several micrometer (Fig. 1a and S1a†). Other G@MC materials also have similar wrinkled multilayer structure (Fig. S2†). After the sulfur incorporation, the stacked structure are preserved (Fig. 1b and S1b†). The TEM images (Fig. 1c and d) show that the G@MC5 material has abundant microporous structure and graphene structure. Thus, the graphene structure stem from the graphene oxides is retained after the activation process, which is advantageous for improving the conductivity of sulfur and discharge products. Furthermore, the microporous structure derived from sucrose is uniformly distributed on the graphene sheets (Fig. 1c and d). Therefore, based on SEM and TEM results, the G@MC5 material has a 3-D structure stacked by graphene sheets and microporous structure. According to the SEM (Fig. 1b) and TEM images (Fig. 1d) of G@MC5–S nanocomposite, no agglomerated sulfur particles or crystalline sulfur are observed on the surface of carbon structure. However, the element mapping (Fig. 1f) reveals that carbon and sulfur are uniformly distributed in the G@MC5–S nanocomposite. Thus, the sulfur is well dispersed in the microporous structure of the G@MC material. Furthermore, the graphene structure also can be seen in the G@MC5–S nanocomposite (Fig. 1e), which is beneficial to the conductivity of G@MC5–S nanocomposite.

Fig. 1 SEM images of G@MC5 material (a) and G@MC5–S nanocomposite (b). TEM images of G@MC5 material (c and d) and G@MC5–S nanocomposite (e). TEM image of G@MC5–S nanocomposite and corresponding elemental mapping images of carbon and sulfur (f).

XRD patterns of the G@MC and G@MC–S materials are shown in the Fig. S3.† The G@MC and G@MC–S materials have broad peaks of carbon, and no obvious sulfur peak is observed in G@MC–S materials (Fig. S3†). This results indicate that sulfur exists in an amorphous state and in a highly dispersed state inside the G@MC materials.^22,24 Raman spectroscopy was used to further investigate the structural features of carbon and sulfur in the G@MC and G@MC–S nanocomposites (Fig. 2). The peaks around 1350 cm⁻¹ (D band) are related to defects and disorder in the carbon material, and the peaks around 1580 cm⁻¹ (G band) are corresponded to the coplanar vibration of sp² hybrid carbon atoms⁴³ in the G@MC and G@MC–S materials. With the increase of the activation agent, the intensity ratio of I_D/I_G increase in the G@MC material (Table S1†). The possible reason is that the activation agent can react with the carbon structure and create detects and disorder in the G@MC material. The crystalline sulfur exhibits a characteristic peak around 475 cm⁻¹, which is related to the A₁ symmetry mode of the S–S bond.⁴³ The G@MC-S nanocomposites do not have the characteristic peak of crystalline sulfur because the sulfur exists in an amorphous state.

Fig. 2 (a) Raman spectra of G@MC materials, (b) Raman spectra of sulfur and G@MC–S materials.

The N₂ adsorption–desorption isotherms were employed to investigate the porous structure of the G@MC materials (Fig. 3a–c). The G@MC2 material exhibits a type I curve in IUPAC classification, indicating the microporous structure of G@MC2 material. With the increase of the ration of KOH versus carbon precursor, the adsorption capacity of N₂ rise up (expect for G@MC6), and the little hysteresis between the adsorption and desorption branches at relative P/P₀ = 0.4–0.6 increase indicating the increasing in mesoporous structure. Thus, the percentage of micropore volume (V_micro/V_T) slightly decrease in the G@MC materials (Table 1) with increase of activation agent. The pore distribution of G@MC materials is calculated by the Density Functional Theory (DFT) method. The G@MC materials have a lot of microporous structure and little mesoporous structure (Fig. 3b and c). All the G@MC materials show three peaks around 0.45, 0.7, and 1.5 nm (Fig. 3b). With increase of the KOH from G@MC2 to G@MC5 material, the peak around 0.4 nm is slightly increase, and the peaks around 0.7 and 1.5 nm are obviously increase, corresponding to the increasing in micropore volume (Table 1). However, in the G@MC6 material, the pore distribution peaks around 0.45 and 0.70 nm decrease, and little mesopore peak is observed (Fig. 3b and c). Thus, the amount of small micropore is decrease and the amount of mesopore is slightly increase with excess KOH. The G@MC5 material has the largest micropore volume (1.84 cm³ g⁻¹) and ultrahigh surface area (3374 m² g⁻¹). Considering the sulfur expansion during the discharge process, the G@MC5 with high pore volume of 2.65 cm³ g⁻¹ can achieve 75 wt% sulfur loading and provide adequate nanospace for sulfur expansion. Thus, both the high percentage of microporous structure and the high sulfur containing have been achieved.

Fig. 3 (a) The N₂ adsorption–desorption isotherms of the G@MC materials and the corresponding pore size distribution curves (b and c) obtained using the DFT method. (d) Thermal analysis curves of G@MC–S nanocomposites.

Table 1 The textural parameters of G@MC materials

Sample	G@MC2	G@MC3	G@MC4	G@MC5	G@MC6
a The specific area (SBET) was calculated by multi-point Brunauer–Emmet–Teller (BET) method.b VT represented the total pore volume (at relative pressure of P/P0 = 0.99).c Micropore volume was determined by applying the Dubinin–Radushkevich (D–R) method to N2 adsorption branch.
SBET^a (m^2 g^−1)	1437	1642	2407	3374	2395
VT^b (cm^3 g^−1)	1.0	1.21	1.78	2.65	2.14
Vmicro^c (cm^3 g^−1)	0.79	0.89	1.3	1.84	1.35
Vmicro/VT	79%	74%	73%	69%	63%

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XPS was used to further investigate the chemical structure of the G@MC–S materials. As shown in the Fig. S4,† G@MC–S materials show typical S 2p_2/3 and S 2p_1/2 peaks with an energy separation of 1.2 eV. The minor peaks at 168.7, 169.1, and 169.2 eV are due to the sulphate species formed by sulfur oxidation in air.³⁸ On the basis of the TG results (Fig. 3d), the sulfur loading are 60.0, 64.6, 67.2, 68.8, 75.4, and 64.4 wt% for the G@MC2–S, G@MC3–S, G@MC4–S, G@MC5–S, G@MC5–S–H, and G@MC6–S nanocomposite, respectively. It is worth noting that the weight loss is between 250 and 400 °C, which is higher than ordinary carbon–sulfur nanocomposite. This phenomenon is attributed to the release of sulfur confined within the micropores of G@MC materials, which needs more driving force and hence higher temperature to overcome the strong capillary force.²⁴

The G@MC–S materials were incorporated into lithium–sulfur batteries to test their electrochemical behavior. The typical voltage capacity profiles of the G@MC–S nanocomposite is shown in Fig. 4a. Two-plateau of sulfur cathode around 2.3 V and 2.1 V is clearly presented, corresponding to the formation of high-order lithium polysulfides (Li₂S_n; where n is typically 4–8) and the sequential reduction of high-order lithium polysulfides into lithium sulfide (Li₂S₂ and Li₂S).⁴ As shown in the Fig. 4b, the G@MC2–S, G@MC3–S, G@MC4–S, G@MC5–S, G@MC5–S–H, and G@MC6–S nanocomposites exhibit the initial discharge capacity of 1245, 1229, 1265, 1235, 1208, and 1249 mA h g⁻¹ at a current rate of 0.2C, and maintain the capacity of 786.6, 788.6, 753.5, 829.5, 788.9, and 713.2 mA h g⁻¹ after 100 cycles at 0.5C, respectively. Furthermore, all the coulombic efficiency of G@MC–S nanocomposites are around 101–102% (Fig. S5†). Considering the sulfur containing in the G@MC–S materials, the performances of the G@MC–S nanocomposites are improved with the increase of the hierarchical microporous structure, pore volume, and surface area. G@MC5–S nanocomposite has the best lithium–sulfur battery performance because of the ultrahigh surface area, large pore volume and high percentage of micropore. The G@MC4 and G@MC6 materials have similar surface area, and the G@MC6 has larger pore volume because of the small amount of mesoporous structure. However, the G@MC4–S nanocomposite has a better lithium–sulfur performance than G@MC6–S nanocomposite because the G@MC4–S nanocomposite has more high percentage of micropore volume (Table 1). Thus, the percentage of microporous structure is important to the performance of lithium–sulfur battery. The long-term cycling performances of the G@MC5–S and G@MC5–S–H nanocomposites at 0.5C after the initial two cycles' activation process at 0.2C are shown in the Fig. 4c. The G@MC5–S and G@CM5–S–H nanocomposites deliver the capacity of 545.3 and 541.3 mA h g⁻¹ after 500 cycles, respectively. Furthermore, all the coulombic efficiency of G@MC–S nanocomposites are around 101% (Fig. 4c).

Fig. 4 (a) Discharge/charge voltage profiles of G@MC5–S nanocomposite, (b) cycle life of G@MC nanocomposites at 0.5C. (c) Long-term cycling performance of G@MC5–S and G@MC5–S–H nanocomposites at 0.5C.

In comparison with the graphene sheets, the G@MC5 material has abundant microporous structure, which is advantageous to restrain the polysulfides diffusion. In comparison with the microporous carbon, the G@MC5 material has sandwich-type graphene and microporous carbon stacked structure, which is advantageous to provide high electrical conductivity for sulfur and discharge products. Thus, the G@MC5–S–H nanocomposite shows excellent lithium–sulfur battery performance, surpassing most previous studies (Table S2†).

A rate capability study was conducted at various rates (0.1, 0.2, 0.5, 1, and 2C) to further investigate the high-rate performance of the G@MC5–S nanocomposite (Fig. 5). When the current rate is increased from 0.1 to 2C, the G@MC5–S nanocomposite maintains the typical two-plateau behavior of the sulfur cathode (Fig. 5a). The sloping plateau during the initial discharge at 1.8 V is attributed to the irreversible reduction of LiNO₃.^44,45 The G@MC5–S nanocomposite acquires a capacity of 785.7 mA h g⁻¹ at a current rate of 2C (Fig. 5b), which is an excellent lithium–sulfur battery performance. Furthermore, the discharge capacity can be mostly recovered when the current density is decrease again from 2 to 0.5C, which shows the significant abuse tolerance of the G@MC5–S nanocomposite.

Fig. 5 The rate capability of the G@MC5–S nanocomposite.

The G@MC5 material can absorb a lot of polysulfides because of the ultrahigh surface area and novel hierarchical microporous structure, as shown in Fig. S6.† The Li₂S₄ was considered as the representative polysulfides, and the DME was used as the solvent. The light yellow Li₂S₄ solution turned almost colorless after the addition of the G@MC material. The excellent adsorption property of the G@MC5 material promotes the remarkable lithium–sulfur battery performance by suppress the diffusion of polysulfides. Furthermore, the underlying reason for the high performance of G@MC5–S nanocomposite was further explored by examining the cathode nanostructure after cycling. The SEM images of the raw G@MC5–S cathode and the G@MC5–S cathode after 500 cycles are shown in the Fig. S7.† The G@MC5–S material maintains the structural integrity after 500 cycles because the G@MC5 material with large pore volume can provide enough nanospace for sulfur expansion.

Our results indicate the rational strategy to fabricate the promising 3-D sandwich-type G@MC nanoarchitectures for lithium–sulfur batteries. As shown in Scheme 1, the rational design of the graphene sheets and microporous carbon stacked structure render the G@MC–S materials with following outstanding advantages: (1) the sandwich-type graphene and microporous structure interlinked network offers high electrical conductivity for the sulfur and discharge products (Li₂S and Li₂S₂); (2) the novel hierarchical microporous structure can absorb and confine the soluble polysulfides because the micropore size of G@MC materials is close to the size of polysulfides,¹² and the 3-D structure stacked by microporous structure and graphene sheets can reduce the random diffusion of the polysulfides; (3) the G@MC materials with abundant micropore and large pore volume can accommodate the volume change of sulfur during cycling and provide sufficient nanospace for the Li₂S₂ and Li₂S deposition, which assures the 3-D sandwich-type framework stacked by graphene sheets and microporous structure; (4) the 3-D sandwich-type structure with abundant microporous structure and a lot of sp² hybrid carbon can facilitate the electrolyte and lithium-ion transportation. Therefore, the G@MC material can serve as an excellent nanoelectrochemical reaction chamber for sulfur cathodes.

Conclusions

In summary, a series of G@MC materials for lithium–sulfur batteries were designed in this study. Both the high pore volume and the high percentage of micropore are necessary for high performance G@MC–S nanocomposite. The 3-D sandwich-type G@MC materials with abundant sp² hybrid carbon can provide 3-D electron transfer pathways and lithium-ion diffusion channels. Furthermore, the hierarchical micropore and graphene stacked structure can absorb and confine the soluble polysulfides. Meanwhile, the high pore volume of G@MC material can provide adequate internal void space for sulfur expansion and ensure the structural integrity of 3-D G@MC during cycling. Thus, the optimized G@MC–S material with high sulfur loading (75.4 wt%) can deliver a capacity of 541.3 mA h g⁻¹ after 500 cycles at rate of 0.5C. In addition, this design strategy is generally applicable for supercapacitors, batteries, catalysis, electrochemical sensors, in which efficient electronic transport and tunable pore structure is critical.

Acknowledgements

This work was supported by the National Natural Science Foundation of China (Grant No. 51225204, 91127044 and U1301244), the National Basic Research Program of China (Grant No. 2013AA050903 and 2012CB932900), and the “Strategic Priority Research Program” of the Chinese Academy of Sciences (Grant No. XDA09010000).

Notes and references

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c5ra24338f

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