Bridging mesoporous carbon particles with carbon nanotubes

Bridging mesoporous carbon particles with carbon nanotubes

Microporous and Mesoporous Materials 98 (2007) 323–329 www.elsevier.com/locate/micromeso Bridging mesoporous carbon particles with carbon nanotubes F...

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Microporous and Mesoporous Materials 98 (2007) 323–329 www.elsevier.com/locate/micromeso

Bridging mesoporous carbon particles with carbon nanotubes Fabing Su a, X.S. Zhao a

a,b,*

, Yong Wang c, Jim Yang Lee

a,c

Department of Chemical and Biomolecular Engineering, National University of Singapore, 4 Engineering Drive 4, Singapore 117576, Singapore b Nanoscience and Nanotechnology Initiative, National University of Singapore, Singapore 117576, Singapore c Singapore-MIT Alliance, National University of Singapore, 4 Engineering Drive 3, Singapore 117576, Singapore Received 4 July 2006; received in revised form 14 September 2006; accepted 21 September 2006 Available online 2 November 2006

Abstract The enhancement of the electrical conductivity (EC) of a porous carbon is highly desirable in many applications, especially in those associated with storage and conversion of electrochemical energy. In this work, we demonstrated an approach to largely increasing the EC of ordered mesoporous carbon (OMC) by bridging the OMC particles with carbon nanotubes (CNTs). Infiltration of the pores of ordered mesoporous SBA-15 silica with a carbon precursor yielded a carbon/mesoporous silica composite, which was further used as a support for Ni catalyst. Subsequently, catalytic growth of CNTs on the Ni-supported composite surface was carried out using the chemical vapor deposition (CVD) method with benzene as the carbon precursor. Removal of the silica framework and the metal catalyst left behind OMC particles bridged with CNTs. The EC of the OMC was increased from 138 S/m (before bridging) to 645 S/m (after bridging). Because of the significant enhancement of EC and the availability of mesopores, the cyclability of the hybrid carbon materials as a negative electrode used in rechargeable lithium-ion batteries was significantly improved. Ó 2006 Elsevier Inc. All rights reserved. Keywords: Mesoporous carbon; Carbon nanotubes; Chemical vapor deposition; Template; Rechargeable lithium ion batteries

1. Introduction Modification of ordered porous carbons (OMCs) prepared by using ordered mesoporous silicas as templates has been shown to improve the carbon performance in various applications [1–3]. OMCs with a high surface area, well-controlled pore structure, and tailorable surface chemistry have been explored in electrochemical energy conversion and storage, such as fuel cells [4–7] and Li-ion batteries [8–10]. The graphitic nature of a carbon is important in terms of its electrical conductivity (EC) and electrochemical activity in fuel cells and rechargeable lithium-ion batteries. The EC of a powdery carbon is determined by the carbon inherent resistance and the contact resistance *

Corresponding author. Address: Department of Chemical and Biomolecular Engineering, National University of Singapore, 4 Engineering Drive 4, Singapore 117576, Singapore. Tel.: +65 65164727; fax: +65 67791936. E-mail address: [email protected] (X.S. Zhao). 1387-1811/$ - see front matter Ó 2006 Elsevier Inc. All rights reserved. doi:10.1016/j.micromeso.2006.09.030

between the carbon particles. A high graphitic nature (crystallinity) that is determined by the length, ordering, and defects of the graphene layers generally implies a low inherent resistance, thus a high EC. While the use of aromatic carbon precursors, such as mesophase pitch [11], acenaphthene [12], poly-vinyl chloride [13], naphthalene [14] and polypyrrole [15,16], in preparation of OMCs has been demonstrated to enhance the EC, such EC enhancements can hardly be further improved due to the small carbon graphene layers confined in the nanospaces of the mesoporous silica templates. Recent studies [13,17] have shown that high temperature annealing and catalytic graphitization can enhance the graphitization degree of OMCs. These strategies, however, have to be implemented at expense of scarifying the porosity of the resultant carbon materials. On the other hand, lowering the contact resistance has been achieved by compression of carbon powdery particles [18] and fabrication of monolithic carbon [19] to increase the effective contact between carbon particles. Here, we demonstrate another approach to lowering the contact

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resistance by bridging carbon particles using carbon nanotubes (CNTs), which are known to possess a high EC. In a recent communication, we reported the preparation of OMCs covered with CNTs using SBA-15 silica with impregnated metal species on both the pore and particle surfaces as template [20]. The metal particles were embedded in the resultant hybrid carbon, which is unwanted in some applications. In this work, we employed an alternative approach to preparing a hybrid carbon similar to that reported in Ref. [20], but without the presence of metal species in the carbon framework. The main difference lies in that instead of using SBA-15 silica with impregnated metal species on both the pore and particle surfaces, SBA-15 silica with impregnated metal species only on its external surface was used as template. As schematically illustrated in Scheme 1, ordered mesoporous SBA-15 silica template was first infiltrated with carbon by using the chemical vapor deposition (CVD) method with benzene as the carbon precursor to yield a carbon/SBA-15 silica composite. Second, a metal nanoparticle catalyst was supported on the surface of the composite. Third, growth of CNTs was carried out on the composite surface. Finally, the silica framework and metal catalyst were removed to leave behind OMC particles bridged with CNTs. The hybrid mesoporous carbon materials displayed a significantly enhanced EC and an improved cyclability when used as a negative electrode in rechargeable lithium-ion batteries.

tube to allow carbon deposition. After 1.5 h, the sample was cooled in pure N2 to obtain the carbon/silica composite. Part of this composite (0.6 g) was mixed with 0.2 g of nickel acetate tetrahydrate (Ni(CH3COO)2 Æ 4H2O, 98%, Alfa Aesar) dissolved in 3 mL of deionized water. The mixture was sonicated for 0.5 h, evaporated, and dried in air at 100 °C for 3 h, followed by calcination at 200 °C for 3 h to obtain a Ni-loaded carbon/silica composite. The growth of CNTs was conducted on the Ni-loaded composite at 900 °C for 0.5 h using the CVD method with benzene vapor as the carbon source. The black sample thus obtained was treated with a 20% HF solution at room temperature and subsequently with a 6 M HNO3 solution at 50 °C for 6 h, followed by washing with copious deionized water and dried in air at 150 °C to obtain a sample denoted as OMC/CNT. For comparison purpose, the other part of the carbon/silica composite without impregnation of Ni catalyst was also used as template for carbon deposition as described above. The carbon sample thus obtained is designed as OMC. For understanding the phase change of the Ni species before and after the CVD process, the Ni-loaded carbon/silica composite was heated at 900 °C for 0.5 h in pure N2 flow (30 cm3/min) without the presence of benzene vapor. The sample thus obtained is denoted NiOx-carbon/silica composite. In addition, a nongraphitizable OMC sample was also prepared by using sucrose as the carbon precursor according to the method reported elsewhere [22].

2. Experimental 2.2. Characterization 2.1. Synthesis of carbon materials Mesoporous silica SBA-15 template was synthesized according to Ref. [21]. The preparation of a carbon/silica composite is described as follows: About 1 g of pure-silica SBA-15 template was placed in a crucible and loaded in a horizontal quartz tube equipped with a furnace. The template was heated to 900 °C with a heating rate of 5 °C/ min in a highly pure N2 flow (30 cm3/min). Subsequently, another N2 flow (30 cm3/min) containing 10 wt% benzene vapor from a liquid bubbler was passed through the quartz

The pore properties of the samples were investigated using physical adsorption of nitrogen at the liquid-nitrogen temperature ( 196 °C) on an automatic volumetric sorption analyzer (Quantachrome, NOVA1200). Prior to measurements, the samples were degassed at 200 °C for 5 h in vacuum. The specific surface areas were determined according to the Brunauer–Emmett–Teller (BET) method in the relative pressure range of 0.05–0.2. The total pore volumes were obtained from the volume of nitrogen adsorbed at the relative pressure of 0.99. The pore size

Scheme 1. Schematic illustration of preparing a hybrid mesoporous carbon (OMC/CNT): (a) SBA-15 silica template, (b) carbon/SBA-15 silica composite after infiltration of the pores of the template with carbon using the CVD method, (c) impregnation of a metal catalyst on the external surface of the carbon/SBA-15 silica composite, (d) growth of CNTs catalyzed by the metal catalyst, (e) OMC/CNT after removal of the silica template and catalyst, (f) the carbon network of OMC/CNT showing the bridging of OMC by CNTs.

F. Su et al. / Microporous and Mesoporous Materials 98 (2007) 323–329

distribution (PSD) curves were derived using the Barrett– Joyner–Halenda (BJH) method from the adsorption branches. The pore sizes were estimated from the maximum positions of the BJH PSD curves. The microscopic features of the samples were observed with a field-emission scanning electron microscope (FESEM) (JSM-6700F, JEOL Japan) operated at 10 kV and a field-emission transmission electron microscopy (FETEM) (JEM 2010F, JEOL, Japan) operated at 200 kV. The mesostructures of the samples were characterized using small-angle X-ray scattering (SAXS) technique on a Bruker NanoStar with CuKa radiation of wavelength k = 0.1542 nm. The graphitic nature of the mesoporous carbon samples was characterized using X-ray diffraction (XRD) technique (XRD-6000, Shimadzu, Japan) with CuKa radiation. Thermogravimetric analysis (TGA) was conducted on a thermogravimetric analyzer TGA 2050 (Thermal Analysis Instruments, USA) with an air flow rate of 100 mL/min and a heating rate of 10 °C/min. The measurement of EC was carried out by pressurizing carbon particles at 6.5 MPa [17]. The direct current (DC) ECs of the samples were obtained from the linear slop of the current–voltage curves [23]. 2.3. Measurement of electrochemical properties The electrochemical properties of the carbon samples as a negative electrode in rechargeable lithium-ion batteries were measured. The working electrode consisted of 80 wt% of the active material, 10 wt% of conductivity agent (carbon black, Super-P) and 10 wt% of binder (polyvinyli-

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dene difluoride, PVDF, Aldrich). Lithium foil was used as the counter and reference electrodes. The electrolyte was 1 M LiPF6 in a 50:50 (w/w) mixture of ethylene carbonate and diethyl carbonate. Cell assembly was carried out in a glove box with the concentrations of moisture and oxygen below 1 ppm. The room-temperature electrode activities were measured using a Maccor-Series-2000 battery tester. The details on the electrode preparation and cell assembly can be found elsewhere [24]. The cells were charged and discharged at a constant current of 0.2 C and the fixed voltage limits were between 3 V and 5 mV. Higher hourly rates (1, 6 and 10 C) were also used for evaluating the performance of the carbon samples. In all tests, the electrodes had the same thickness and were loaded to the tester according to the same procedure. 3. Results and discussion 3.1. Macroscopic observations Fig. 1 shows the FESEM images of NiOx-carbon/SBA15 silica composite, OMC/CNT and CNTs bridging the OMC particles. It can be seen that the morphology of the OMC composed of bamboo-shaped primary particles is similar to that of the SBA-15 silica template (see Figure S1-a in Supporting Information (SI)). Conglutination of the OMC particles by non-templated carbon as observed on the sample prepared using the sucrose-impregnation method [22] (see Figure S1-b in SI) has been avoided by using the CVD method (Figure S1-c in SI). Fig. 1a also reveals a uniform dispersion of the NiOx nanoparticles

Fig. 1. FESEM images of (a) NiO-carbon/silica composite, (b,c) sample OMC/CNT of different magnifications, and (d) CNTs bridging the OMC particles.

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on the surface of the carbon/silica composite. The diameter of a CNT is often related to the particle size of a metal catalyst on which it grows [25]. Fig. 1b clearly shows a large number of CNTs with a length of several micrometers covering the surfaces of the OMC particles, indicating that the CNTs grew catalytically on the external surface of the OMC bamboo-shaped primary particles. White dots (Ni particles) on the tips of the CNTs can be more clearly observed from Figure S2 (see SI), suggesting the growth

of the CNTs followed the tip-growth model [25]. As can be seen from Fig. 1c, these CNTs bridge the OMC particles to form a carbon network. Fig. 1d shows the morphology of the CNTs with a uniform diameter of around 30–40 nm. Fig. 2 compares the different behaviors of samples OMC (the left-hand sample) and OMC/CNT (the right-hand sample) dispersed in ethanol by shaking. It was observed that after 6 s, sample OMC/CNT began to settle down (Fig. 2a). After 30 s, the sedimentation of sample OMC/ CNT became significant (Fig. 2b). After 1200 s, most of OMC/CNT particles settled down to the bottom of the container while sedimentation of OMC particles just began (Fig. 2c). The much faster sedimentation rate of the OMC/ CNT particles than that of the OMC particles is due to the bridging effect of the OMC particles by the CNTs in sample OMC/CNT, which is very similar to bridging flocculation of colloidal particles by polymers. Fig. 2d shows that the OMC/CNT sample displays magnetic properties because of the presence of residual Ni metal as will be confirmed by the XRD and TG data below.

3.2. TEM observations

Fig. 2. Photographs of samples OMC (the left-hand one) and OMC/CNT (the right-hand one) dispersed in ethanol by shaking, followed by letting the particles settle down for (a) 6 s, (b) 30 s, (c) 1200 s, and photographs of OMC/CNT with the magnetic property (d).

The microscopic features of samples OMC and OMC/ CNT are presented in Fig. 3. Fig. 3a and b shows the TEM images of sample OMC taken along the [0 0 1] and [1 0 0] directions, respectively. The interplanar distance between the arrays was estimated to be about 10 nm, consistent with that of the SBA-15 silica template (see Figure S3 in SI). Fig. 3c reveals the presence of numerous CNTs bridging the OMC particles in sample OMC/CNT. The diameters of these CNTs are in the range of 30–40 nm.

Fig. 3. TEM images of OMC (0 0 1) direction (a) and (1 0 0) direction (b), OMC/CNT (c), and a CNT (d).

F. Su et al. / Microporous and Mesoporous Materials 98 (2007) 323–329

3.3. SAXS patterns

(100)

(110)

a Intensity (a.u)

(*) carbon (+) Ni (o) NiO

*+ +

* o

*

o

+

c

+

o

+

b o a

20

30

40 50 60 2 Theta (Degree)

70

80

Fig. 5. XRD patterns: (a) OMC, (b) NiO-carbon/silica composite, and (c) OMC/CNT.

3.4. XRD patterns

Fig. 4 shows the SAXS patterns of SBA-15 silica, OMC and OMC/CNT. All samples display well-resolved (1 0 0) (1 1 0), and (2 0 0) peaks, demonstrating the presence of two-dimensionally ordered hexagonally arranged mesopore arrays. The position of the (1 0 0) peak observed at about 0.83 degree two theta on sample SBA-15 is seen at about 0.92 degree two theta on samples OMC and OMC/ CNT due to the structural shrinkage of the template upon high-temperature thermal treatment. The interplanar distances d1 0 0 of samples OMC and OMC/CNT were calculated to be about 9.6 nm, consistent with the TEM observations. In addition, the SAXS pattern of sample OMC shows the presence of the well-resolved (1 1 0) and (2 0 0) peaks with a strong intensity, indicating a long-range structural ordering. Upon bridging by CNTs, the pore structure of the OMC was not altered in spite of the lowered intensities of the (1 0 0) (1 1 0), and (2 0 0) peaks.

X3

(200)

b X3

a X3

b c

X3

c 0.5

*

Intensity (a.u)

Fig. 3d shows the fringe-lattice of one of the CNTs with a diameter of around 30 nm. The wall thickness and inner diameter of this CNT are about 9 and 12 nm, respectively. It can also be seen that the straight graphene layers of multiwalls are parallel to the axis of the CNT. These graphene layers with an interlayer distance of about 0.34 nm can be clearly seen from the inset in Fig. 3d, implying a high crystallinity of the CNTs bridging the OMC particles of sample OMC/CNT. Such highly graphitic CNTs will facilitate the electronic transport along the axis of the CNTs [26]. It should be noted that in comparison with the hybrid carbon material prepared using the method described in Ref. [20], the present method allowed us to prepare hybrid carbon materials with highly ordered mesopores without the presence of metal particles embedded in the carbon framework (a comparison of Fig. 3a with Fig. 1d of Ref. [20] revealed this conclusion).

327

1.0 1.5 2 Theta (Degree)

2.0

2.5

Fig. 4. SAXS patterns: (a) OMC/CNT, (b) OMC, and (c) SBA-15 silica.

The XRD patterns shown in Fig. 5 demonstrate that sample OMC prepared using the CVD method with benzene as the carbon precursor exhibits a broad peak at around 25° 2h (Fig. 5a). This peak corresponds to the (0 0 2) diffraction of graphitic carbon, suggesting a low graphitization degree of the sample. The XRD pattern of the NiO-carbon/silica composite shown in Fig. 5b displays characteristic peaks at 37.3°, 43.3°, 62.9°, and 75.4° 2h, which can be indexed according to the (1 0 1) (0 1 2) (1 1 0) and (1 1 3) diffractions of NiO [27]. The two small peaks at 44.5° and 51.9° 2h may be due to the presence of minor Ni metal formed by reduction of NiO by carbon species at high-temperatures. Here, NiOx means a mixture of NiO and Ni. The strong diffraction peak with a high intensity at about 26.2° 2h together with a small peak at about 42.8° 2h observed on sample OMC/CNT (Fig. 5c) are believed to stem from the CNTs. In addition, the other three peaks observed at about 44.7°, 52.0°, and 76.6° 2h can be indexed according to a face-centered cubic structure of crystalline Ni corresponding to the (1 1 1) (2 0 0) and (2 2 0) reflections, respectively (JCPDS file No. 4-485), again showing the presence of Ni metal in sample OMC/ CNT. It is noted that the strong peak at about 44.7° 2h is not only due to the (1 1 1) reflection of Ni metal, but also the (1 0 1) reflection of graphitic carbon. No NiO diffraction peaks can be found from Fig. 5c (especially at around 37.3°), suggesting a complete phase transformation of NiO to Ni during the CVD process because of the reduction of NiO to metal Ni by hydrogen from the dehydrogenation of benzene [28]. It is believed that it was actually the Ni metal nanoparticles that served as the catalytic sites for the growth of the CNTs of sample OMC/CNT [25]. Since the growth of the carbon nanotubes (CNTs) in this work followed the tip-growth model, Ni particles are buried or encapsulated within the graphene layers at the top of CNTs and can not be completely removed by acid treatment.

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3.6. Nitrogen adsorption Fig. 7 shows the N2 adsorption–desorption isotherms and BJH PSD curves of samples OMC and OMC/CNT. It is seen that the isotherms of both samples can be classified as the Type-IV isotherm with a H2 hysteresis loop, indicating they are mesoporous materials. The pore sizes of samples OMC and OMC/CNT are centered at about 3.5 and 3.6 nm, respectively, indicating the high-temperature CNT growing process did not alter the mesopore structure of the OMC sample. The surface areas and pore 100

OMC OMC/CNT

3.6 nm

1

10 Pore size (nm)

0.2 0.4 0.6 0.8 Relative pressure (P/P0 )

100

1.0

Fig. 7. Adsorption–desorption isotherms and BJH–PSD curves (inset) of synthesized carbons OMC and OMC/CNT.

volumes of samples OMC and OMC/CNT were calculated to be about 654 and 484 m2/g, and 0.71 and 0.53 cm3/g, respectively. The lower surface area and pore volume of sample OMC/CNT than that of OMC are due to the presence of dense CNTs in sample OMC/CNT. 3.7. Electrochemical properties in lithium-ion batteries Fig. 8 compares the cycling performances of OMC and OMC/CNT used as a negative electrode in lithium ion batteries measured at a constant current of 0.2 C and fixed voltage windows of 5 mV to 3 V. The initial capacities of OMC and OMC/CNT were 755 and 680 mA h/g, respectively, substantially higher than the theoretical capacity of graphite (372 mA h/g). The higher initial capacity of OMC than that of OMC/CNT may be attributed to its more nanopores and/ or low graphitization degree. After 45 cycles of discharging and charging, about 87.5% of the initial capacity of sample OMC/CNT was maintained. However, only about 71.1% of the initial capacity of sample OMC was observed after the 45 cycles. Thus, improvement of cyclability by bridging the OMC particles with CNTs is obvious (about 0.27% loss

Specific capacity (mAh/g)

20 0

3.5 nm

800 Deriv.Weight (%/ oC)

Weight (%)

40

OMC OMC/CNT

200

0 0.0

80 60

400

3

Fig. 6 shows the thermogravimetric behaviors of samples OMC and OMC/CNT. The negligible residue above the temperature of about 800 °C of sample OMC shows the complete removal of the silica template by aqueous HF solution. Thus, the residue of sample OMC/CNT can be considered as the mass of NiO (6.0 wt%), which were converted to about 4.7 wt% of Ni metal in sample OMC/ CNT. This is supported by the EDX data (not shown here), from which it was estimated that the residual Ni metal in sample OMC/CNT was about 5.6 wt%. The TGA curves are rather interesting. First, only one relatively broad weight-loss peak centered at about 650 °C is seen on sample OMC whereas two sharp peaks centered at about 550 and 630 °C are seen on sample OMC/CNT. The peak of sample OMC was due to the combustion of the templated carbon framework. This combustion temperature of the carbon framework of sample OMC/CNT peak was significantly lowered to about 550 °C due to the catalysis of the NiO derived from Ni residual in air. The peak at about 630 °C of sample OMC/CNT was due to the combustion of the CNTs bridging the OMC particles, again under the catalysis on the NiO. A simple calculation [16] tentatively showed that sample OMC/CNT contains approximately 70 wt% of OMC, 25 wt% of CNTs and 5 wt% of Ni.

Adsorbed volume (cm /g, STP)

3.5. Thermogravimetric analysis

200 400 600 800 Temperature (oC)

200

400 600 Temperature (oC)

800

Fig. 6. Thermogravimetric analysis (TG) and DTG (inset) of OMC and OMC/CNT.

600

400 OMC/CNT OMC

200

0 0

10

20 30 Cycle number

40

50

Fig. 8. Cycling performances of OMC and OMC/CNT anodes.

F. Su et al. / Microporous and Mesoporous Materials 98 (2007) 323–329 Table 1 The specific capacities of the carbon samples at various specific currents Specific current (mA/g)

0.2 C

1C

6C

10 C

OMC OMC/CNT MCMB

755 680 330

686 625 150

614 487 32

252 354 /

per cycle for sample OMC/CNT against about 0.64% loss per cycle for sample OMC after 45 cycles of discharging and charging). Such an improvement can be immediately attributed to the contribution of the highly crystalline CNTs bridging the OMC particles, thus enhancing the EC. It should be noted that the ability of storing a large amount of lithium ions of CNTs [29] may have also contributed to the observed improvement on cyclability. Table 1 reports the rate capabilities of the two carbon sample, together with that of mesophase carbon microbeads (MCMB), which is taken as the benchmark in the rechargeable Li-ion battery community [30]. The initial specific capacities of 755 and 680 mA h/g of samples OMC and OMC/CNT obtained at the constant current of 0.2 C were decreased to 614 and 487 mA h/g when the current rate was 6 C. Nevertheless, the initial specific capacities at 6 C are significantly higher than that of the MCMB, which gave a capacity of about 32 mA h/g at the same current rate (6 C). Thus suggests the great rate performance of ordered mesoporous carbon as a negative electrode used in Li-ion batteries. When the current rate was increased to 10 C, the initial specific capacity of OMC (252 mA h/g) was lower than that of OMC/CNT (354 mA h/g). During a high-current-rate electrochemical process, the EC of an electrode material is a key factor determining the specific capacity. Measurement of EC showed that sample OMC/CNT displayed an EC of 645 S/m, which is much higher than that of sample OMC (138 S/m), as well as higher than that of similar mesoporous carbon materials reported in the literature [14,19]. The profound increase in EC is interpreted to be because of the presence of CNTs in sample OMC/CNT, which bridged the OMC particles, thus lowering the contact resistance and in the meantime facilitating electron transport.

4. Conclusions In summary, we have demonstrated that the electrical conductivity of OMCs can be significantly enhanced from 138 to 645 S/m by bridging mesoporous carbon particles with CNTs. This was achieved by growing CNTs on the external surface of a SBA-15 silica/carbon composite, followed by removal of the silica framework. The method described in this work allowed us to control the growth of CNTs while preserving the mesoporous carbon structure. Because of the significant enhancement of EC and the presence of mesopores, the cyclability of the hybrid carbon materials as a negative electrode in rechargeable lithium-ion batteries has been greatly improved in comparison with that of OMC materials without CNTs (a 0.27%

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loss per cycle against a 0.64% loss per cycle). Additionally, a better rate performance of the OMC/CNT hybrid carbon at a high rate (10 C) was observed. The preparation method for such hybrid carbon materials described in this work offers an approach to enhancing the EC of carbons for applications associated with storage and conversion of electrochemical energy. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.micromeso. 2006.09.030. References [1] M. Choi, R. Ryoo, Nat. Mater. 2 (2003) 473. [2] Z. Guo, G. Zhu, B. Gao, D. Zhang, G. Tian, Y. Chen, W. Zhang, S. Qiu, Carbon 43 (2005) 2344. [3] Z. Li, S. Dai, Chem. Mater. 17 (2005) 1717. [4] S.H. Joo, S.J. Choi, I. Oh, J. Kwak, Z. Liu, O. Terasaki, R. Ryoo, Nature 412 (2001) 169. [5] W.C. Choi, S.I. Woo, M.K. Jeon, J.M. Sohn, M.R. Kim, H.J. Jeon, Adv. Mater. 17 (2005) 446. [6] J. Ding, K.-Y. Chan, J. Rena, F. Xiao, Electrochim. Acta 50 (2005) 3131. [7] F. Su, J. Zeng, X. Bao, Y. Yu, J.Y. Lee, X.S. Zhao, Chem. Mater. 17 (2005) 3960. [8] H. Zhou, S. Zhu, M. Hibino, I. Honma, M. Ichihara, Adv. Mater. 15 (2003) 2107. [9] J. Fan, T. Wang, C. Yu, B. Tu, Z. Jiang, D. Zhao, Adv. Mater. 16 (2004) 1432. [10] I. Grigoriants, L. Sominski, H. Li, I. Ifargan, D. Aurbach, A. Gedanken, Chem. Commun. (2005) 921. [11] H. Yang, Y. Yan, Y. Liu, F. Zhang, R. Zhang, Y. Meng, M. Li, S. Xie, B. Tu, D. Zhao, J. Phys. Chem. B 108 (2004) 17320. [12] T.W. Kim, I.S. Park, R. Ryoo, Angew. Chem. Int. Ed. 42 (2003) 4375. [13] A.B. Fuertes, S. Alvarez, Carbon 42 (2004) 3049. [14] C.H. Kim, D.K. Lee, T.J. Pinnavaia, Langmuir 20 (2004) 5157. [15] C.M. Yang, C. Weidenthaler, B. Spliethoff, M. Mayanna, F. Schu¨th, Chem. Mater. 17 (2005) 355. [16] A.B. Fuertes, T.A. Centeno, J. Mater. Chem. 15 (2005) 1079. [17] M. Sevilla, A.B. Fuertes, Carbon 44 (2006) 468. [18] J. Sa´nchez-Gonza´lez, A. Macı´as-Garcı´a, M.F. Alexandre-Franco, V. Go´mez-Serrano, Carbon 43 (2005) 741. [19] L. Wang, S. Lin, K. Lin, C. Yin, D. Liang, Y. Di, P. Fan, D. Jiang, F.-S. Xiao, Micropor. Mesopor. Mater. 85 (2005) 136. [20] F. Su, X. Li, L. Lv, X.S. Zhao, Carbon 44 (2006) 801. [21] D. Zhao, J. Feng, Q. Huo, N. Melosh, G.H. Fredrickson, B.F. Chmelka, G.D. Stucky, Science 279 (1998) 548. [22] S. Jun, S.H. Joo, R. Ryoo, M. Kruk, M. Jaroniec, Z. Liu, T. Ohsuna, O.J. Terasaki, J. Am. Chem. Soc. 122 (2000) 10712. [23] M. Han, S.H. Chan, S.P. Jiang, J. Power Sources 159 (2006) 1005. [24] Y. Wang, J.Y. Lee, T.C. Deivaraj, J. Electrochem. Soc. 151 (2004) A1804. [25] K. Otsuka, Y. Abe, N. Kanai, Y. Kobayashi, S. Takenaka, E. Tanabe, Carbon 42 (2004) 727. [26] A. Javey, J. Guo, Q. Wang, M. Lundstrom, H.J. Dai, Nature 424 (2003) 654. [27] M. Bououdina, D. Grant, G. Walker, Carbon 43 (2005) 1286. [28] Y. Tian, Z. Hu, Y. Yang, X. Wang, X. Chen, H. Xu, Q. Wu, W. Ji, Y. Chen, J. Am. Chem. Soc. 126 (2004) 1180. [29] M. Terrones, Annu. Rev. Mater. Res. 33 (2003) 419. [30] H. Wang, T. Abe, S. Maruyama, Y. Iriyama, Z. Ogumi, K. Yoshikawa, Adv. Mater. 17 (2005) 2857.