Structural evolution of carbide-derived carbons upon vacuum annealing

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Structural evolution of carbide-derived carbons upon vacuum annealing Sebastian Osswald

a,* ,

John Chmiola b, Yury Gogotsi

c

a

Department of Physics, Naval Postgraduate School, 1 University Circle, Monterey, CA 93943, USA Lawrence Berkeley National Laboratory, 1 Cyclotron Road, Berkeley, CA 94122, USA c Department of Materials Science and Engineering, and A.J. Drexel Nanotechnology Institute, 3141 Chestnut Street, Philadelphia, PA 19104, USA b

A R T I C L E I N F O

A B S T R A C T

Article history:

Microstructure and surface moieties of porous carbons play a significant role in affecting

Received 27 March 2012

their performance in a variety of applications. While it is well known that high-tempera-

Accepted 8 June 2012

ture treatments of porous carbons can influence the microstructure, no systematic studies

Available online 17 June 2012

have been done on carbide-derived carbons. We show that vacuum annealing increases the pore volume and specific surface area of titanium carbide-derived carbon with no significant change in the pore size up to 1500 C. This treatment produces porous carbons with subnanometer porosity and a specific surface area up to 2000 m2/g, while treating the samples at temperatures above 1600 C increases the pore size above 1 nm because of graphitization and collapse of the micropore structure. The results demonstrate that vacuum treatment can be used to further tune the pore structure and potentially the surface functionality of carbide-derived carbons for supercapacitor electrodes, gas chromatography, sorption, sensing and other applications. Vacuum annealing of carbide-derived carbon is therefore a suitable alternative to conventional microstructure modification methods, such as gas or liquid phase activation, which are subject to substantial sample loss and result in additional surface functionalization. Published by Elsevier Ltd.

1.

Introduction

Porous amorphous carbons such as activated [1] or carbidederived carbons (CDC) [2] are widely used in applications ranging from water filtration and capacitive water deionization [3] to chemical sensing, supercapacitor electrodes, and gas sorption and storage [4]. Pore size, specific surface area, surface chemistry, and level of graphitization are the key parameters that determine the performance of porous carbons in these applications. In most cases, microstructure and porosity control are limited during synthesis, requiring post treatments, such as activation, to further increase surface area and pore volume. Carbons produced by chemical

* Corresponding author. E-mail address: [email protected] (S. Osswald). 0008-6223/$ - see front matter Published by Elsevier Ltd. http://dx.doi.org/10.1016/j.carbon.2012.06.016

or physical activation have oxygenated surfaces and contain a large variety of functional groups, such as carbonyl, carboxyl, hydroxyl and others, which render the carbon surface hydrophilic [1]. While the hydrophilic nature of the carbon surface is beneficial for some applications, such as water desalination [3], this surface functionalization is undesirable in supercapacitor electrodes, where CDC is widely used [2], because electrolytes based on organic solvents and ionic liquids are incompatible with water. Drying the carbon materials prior to cell assembly may be responsible for up to 30% cost of supercapacitors [5]. However, even if extensive drying is utilized, the oxygen- and hydrogen-containing functional groups also directly alter the

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electrochemical properties and may lead to various unwanted side reactions [6–9]. Most CDCs have a lower oxygen content compared to activated carbons as they are synthesized in an oxygen-free atmosphere [7,8]. However, when the material is exposed to the environment after synthesis, all dangling bonds are quickly saturated with oxygen [12], especially for materials synthesized below 1000 C. Moreover, to achieve a larger pore volume, CDC has been activated [13,14] and activation during synthesis has been proposed recently [15]. Finally, in addition to unfavorable surface functionalization, activation also leads to significant sample loss, particularly when synthesizing high surface area CDCs. Similarly, high temperature treatments under high-vacuum conditions, commonly referred to as vacuum annealing, can be used to modify the microstructure of carbonaceous materials. While vacuum annealing may allow for similar improvements in surface area and control of porosity, it does not suffer from sample loss and does result in unfavorable surface functionalization. To the best of our knowledge, the effect of annealing in vacuum on the physical and chemical characteristics of CDC has only been studied below 1000 C, a temperature range insufficient for major structural modification [16]. In this study we report on the high temperature (up to 2000 C) vacuum annealing of CDC, and demonstrate that vacuum annealing is suitable method to control the microstructure and pore size of CDC, while removing oxygen and other undesirable functional groups upon high-temperature treatment.

2.

Experimental

Titanium carbide (TiC) CDC chlorinated at 600 C for 3 h, followed by a 90 min treatment in H2 at 600 C, herein referred to as TiC-CDC, was purchased from Skeleton Technologies Ltd. (Tartu, Estonia). The as-received TiC-CDC powder was loaded into graphite crucibles, placed into a resistive graphite element diffusion pump backed furnace (Vacuum Atmospheres) and treated at different temperatures in the range 1000–2000 C in vacuum (106 mmHg). To prevent oxidation of the samples at high temperature, the furnace was held for 2 h at 200 C under high vacuum conditions (<104 Torr) to outgas any trapped gasses. The samples were then heated at 10 C/min and held for 2 h at the desired temperature, ranging from 1000 to 2000 C. After annealing, the samples were cooled down under vacuum at the same rate. Changes in structure and porosity of TiC-CDC upon high temperature vacuum-annealing were studied using a combination of Raman spectroscopy, X-ray diffraction (XRD), gas porosimetry, and high-resolution transmission electron microscopy (HRTEM). Raman analysis was performed using a Renishaw 1000/2000 Raman micro spectrometer with a 514 nm Ar+ laser (1800 l/mm grating, 50· Objective, max. 800 W/cm2) in backscattering geometry. Data analysis was performed using GRAMS-32 and WiRE 2.0 software from Renishaw. XRD analysis was performed using a Siemens D500 powder diffractometer in 2H-configuration, equipped with a ˚ ) and a 1500 W copper fine focus tube (CuKa, k = 1.54056 A graphite monochromator. A step size of 0.02 (2H) and a col-

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lection time between 1 and 5 s per step were chosen for the analysis. Diffraction patterns were analyzed using Jade+(MDI) analysis software and Jade MDI powder diffraction library. JEOL 2010F field emission TEM operating at accelerating voltage of 100 and 200 kV was used for high-resolution imaging. HRTEM samples were prepared by dispersing the CDC samples in isopropyl alcohol over a copper grid coated with a lacey carbon film. Porosity was studied using Argon adsorption at 77 K and the nonlocal density functional (NLDFT) theory kernel [17] for slit shaped pores provided by Quantachrome Instruments. Samples were outgassed at 300 C overnight prior to analysis. To ensure accuracy, a minimum sample mass of 20 mg was used, which required longer equilibration times, but minimized potential variation and inaccuracy.

3.

Results and discussion

3.1. High-resolution transmission electron microscopy and X-ray diffraction Representative HRTEM micrographs for the as-produced TiCCDC (Fig. 1a) and samples treated at 1400 C (Fig. 1b) and 2000 C (Fig. 1c) show an increase in ordering as the treatment temperature increases. As-produced TiC-CDC exhibits a highly amorphous structure (Fig. 1a) as discussed in detail in previous studies (see Refs. [12,14]). As the vacuum-annealing temperature is increased to 1400 C (Fig. 1b), the appearance of parallel but curved fringes indicates the onset of graphitization, but no planar graphitic carbon is formed. Upon further increase in annealing temperature to about 2000 C (Fig. 1c), the fringes straighten and grow in both length and number of layers, suggesting a further increase in the level of graphitization and the formation of graphitic ribbons. Results from XRD analysis of the TiC-CDC samples are consistent with TEM data. The XRD pattern of as-produced TiC-CDC and samples annealed at temperatures below 1400 C do not display any distinct diffraction features. At 1600 C, straightening and stacking of fringes upon graphitization lead to the appearance of the broad (0 0 2) diffraction peak indicating an increase in ordering in the direction perpendicular to the graphitic planes. With a further increase in annealing temperature to about 2000 C, the (0 0 2) peak fully develops, suggesting the presence of planar graphite nanoribbons with a well-ordered layer structure. This does not happen during the TiC-CDC synthesis, even at 1200 C. The maximum of (0 0 2) is around 2H = 26.2, which corresponds to an interplanar lattice spacing ˚ . The (0 0 2) peak is asymmetric and shifted toof d0 0 2 = 3.40 A wards lower scattering angles as compared to bulk graphite ˚ ). The asymmetric line shape is the (2H  26.7, d0 0 2 = 3.35 A result from the simultaneous contributions from species with varying degrees of structural ordering, as previously reported for other disordered carbon materials at the onset of graphitization [18].

3.2.

Raman spectroscopy

Raman spectroscopy is a powerful characterization tool for carbon allotropes and can provide deep insights into the

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Fig. 1 – Structural development of CDC during vacuum-annealing (a–c) TEM micrographs showing structure development with increasing annealing temperature. (d) Evolution of the (0 0 2) XRD peak of graphite with temperature.

structural changes that take place upon vacuum annealing [19]. The Raman spectrum of amorphous carbon depends strongly on the relative amounts of sp3, sp2, and sp1-bonded sites [20]. In particular, clustering and orientation of the sp2 sites are crucial parameters, since the level of graphitization determines most of the physical properties of carbon materials with high sp2 contents, such as CDC. Fig. 2a shows the first order Raman spectra of TiC-CDC vacuum-annealed at different temperatures in comparison to as-received TiC-CDC and bulk graphite (single crystal). To account for the differences in peak shape, number of peaks, and intensity distribution, two different fitting procedures were used for highly amorphous, as-produced TiC-CDC (Fig. 2b) and highly graphitized TiC-CDC (Fig. 2c), respectively. Like other disordered carbon materials, TiC-CDC exhibits two broad Raman peaks centered around 1360 and 1600 cm1, corresponding to the disorder-induced D band, associated with a double-resonance process (inter-valley scattering), and the G band of graphitic carbon, assigned to the in-plane stretching vibrations of sp2 sites, respectively [21,22]. A thorough analysis of G band of the asproduced TiC-CDC (Fig. 2b) reveals that this peak results from the overlap of two separate Raman features, herein referred to as G1 and G2, which can be fitted using two peaks centered at 1550 and 1595 cm1, respectively. While the Raman band at 1595 cm1 (G2) is assigned to graphitic sp2 carbon, the peak around 1550 cm1 (G1) is believed to originate from the G band contributions of highly amorphous sp2 carbon phases

Fig. 2 – Raman spectra of TiC-CDC vacuum-annealed at different temperatures in comparison to single crystal, bulk graphite (a). Raman analysis and peak fitting for asproduced (b) and vacuum-annealed (2000 C) TiC-CDC (c).

[22,23]. Similarly, the D band in the Raman spectra of CDC, particularly for CDCs synthesized at low chlorination temperatures, can be fitted as an overlap resulting from the simultaneous contributions of graphitic (1360 cm1) and amorphous (1170 cm1) sp2 sites [22–24]. However, in this study we will not distinguish between the different D band contributions and will use a single peak for data analysis (spectral fitting). At higher annealing temperatures (>1400 C), D and G bands become narrow and a third Raman feature, commonly referred to as D 0, is observed around 1620 cm1 (Fig. 2(c)). The D 0 mode, which does not exist in pure graphite, is a double-resonant Raman feature (intra-valley scattering) and has been assigned to the in-plane vibrations of the outer parts of graphitic domains [25,26]. All spectra were normalized with respect to the maximum G band intensity. It should be noted that while Raman spectra obtained for low (1000 C) and high (2000 C) annealing temperatures can be well fit using the procedures in Fig. 2b and Fig. 2c, respectively, analysis of Raman data recorded at intermediate annealing temperatures is more complex due to the overlap of the various Raman features. Fig. 3 shows the changes in frequency and full width at half maximum (FWHM) of the D, G, and D 0 Raman bands with increasing annealing temperature. Full and open symbols represent Raman data obtained using the fitting procedures shown in Fig. 2b (without D 0 band) and Fig. 2c (with D 0 band), respectively. The markings along the y-axes indicate the values measured for as-received TiC-CDC produced at 600 C. The changes in the position of the G1 and G2 Raman bands are displayed in Fig. 3a. The as-received TiC-CDC 600 C exhibits a G2 frequency of 1596 cm1, which is slightly lower than the 1600 cm1 recorded after vacuum annealing at 1000 C. Samples annealed at 1200 and 1400 C yield similar values and do not exhibit noticeable changes in the G2 frequency. A further increase in annealing temperature, from 1400 to 1800 C, results in a steady downshift of the G2 band to 1585 cm1. Similarly, the G1 band Raman shift remains fairly constant between 1000 and 1600 C (around 1540 cm1), but decreases with further increase in annealing temperature to 1510 cm1 at 2000 C. With a value of 1550 cm1, as-received TiC-CDC 600 C showed the highest G1 Raman shift of all samples. Similar but less pronounced trends were observed when accounting for the D’ band contributions (open symbols) during peak fitting. The double resonant D and D 0 bands both reveal a downshift in Raman frequency with increasing annealing temperature (Fig. 3b). The D band shifts from 1348 cm1 at 1000 C to

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Fig. 3 – Changes in frequency (a and b) and FWHM (c and d) of the D, G, and D 0 Raman bands after high-temperature vacuum annealing of TiC-CDC. Open and closed symbols represent data obtained with and without D 0 band contributions, respectively. The ‘‘X’’ indicates the values of as-received TiC-CDC.

1342 cm1 at 2000 C, whereas the D 0 band decreases from 1620 to 1615 cm-1. Fig. 3c and d show the corresponding changes in FWHM for the G1 and G2, and D and D 0 Raman bands, respectively. The FWHM of the high-frequency component of the G band, G2, decreases with increasing annealing temperature from 76 cm1 (1000 C) to 65 cm-1 (1600 C), but exhibits no further change at 1800 and 2000 C (Fig. 2c). The FWHM of G1 shows only small changes between 1000 C (144 cm1) and 1600 C (130 cm1), but decreases sharply for higher annealing temperatures, reaching a value of 85 cm1at 1800 C. Similarly to G2, the FWHM of the D band decreases between 1000 C (225 cm1) and 1600 C (77 cm1) and exhibits only small changes with further increase in annealing temperature to 1800 C (60 cm1) and 2000 C (77 cm1). The FWHM values of D 0 band range between 46 and 48 cm1 at 1400, 1600, and 1800 C, but decrease to about 40 cm1 after annealing at 2000 C. Fig. 4 displays the ratios between the intensity (peak height) and integrated intensity (peak area) of the D and G2 bands. The D/G2 intensity ratio slightly decreases from 1.48 at 1000 C to 1.34 at 1400 C, followed by an increase to 1.65 at 1800 C. At high annealing temperatures (2000 C) the D/G2 intensity ratio decreases again, reaching a final value of 1.37. In contrast, the integrated D/G2 intensity ratio decreases from 6.1 at 1000 C to 1.7 at 2000 C, with the primary decrease occurring between 1200 C (5.8) and 1600 C (2.4). Similar trends were observed when accounting for the D 0 band contributions (open symbols). The graphitization process and the related changes in the Raman spectra upon annealing can be explained using the three-stage model developed by Ferrari and Robertson [21,23]. Due to the high content of sp2-carbon in CDC (>80%) [27], graphitization of TiC-CDC 600 C by high temperature

Fig. 4 – Intensity ratio (blue squares) and integrated intensity ratio (red circles) between D and G2 Raman bands. Open and closed symbols represent data obtained with and without D 0 band contributions, respectively. The X indicates the values of as-received TiC-CDC.

vacuum annealing is considered a reverse stage 2 process, consisting of a transition from amorphous sp2-carbon to nanocrystalline graphite. For a better understanding, it is appropriate to look at the molecular interpretation of the D and G bands. The G peak results from bond stretching of all pairs of sp2 atoms in both, rings and chains. The D peak is due to the breathing modes of the aromatic sp2 rings, but is symmetry forbidden in defect-free bulk graphite and only Raman active in the presents of defects and disorder. [23,24] The intensity (peak height) is proportional to the number of sixfold sp2 carbon rings. The width of the D band (FWHM) is related to the ordering and size distribution of sp2 rings. Presence of five-, seven-, and eightfold rings in CDC [28] leads to an increase of the FWHM. Below 1400 C, both G1 and G2 do not reveal significant changes in Raman shifts, indicating that the sp2-carbon content remains constant with increasing annealing temperature.

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At the same time, the D band frequency slightly decreases, which may be explained by a decline in the number of smaller aromatic rings [29]. This assumption is also supported by the decreasing FWHM of the D band over the same temperature range. However, the total frequency shift is small, and an accurate interpretation of changes in the D band position may be hindered by the overlapping contributions of other Raman features not accounted for in this study [23]. The D/G1 intensity ratio is another measure used for evaluating the level of graphitization in sp2-carbons. Fig. 3 shows that the D/G2 intensity ratio remains almost constant between 1000 and 1400 C, suggesting that both the degree of layer stacking and the number of sixfold rings remain largely unchanged and that no graphitization takes place, in good agreement with XRD data. While TEM analysis revealed some increase in ordering around 1400 C, these changes are more related to 2-D crystalline growth and not layer growth and stacking and therefore not seen as graphitization using Raman. At higher annealing temperatures, between 1400 and 1800 C, several changes in the Raman spectra are observed. G1 and G2 exhibit a decrease in both Raman frequency and FWHM while the D/G2 intensity ratio increases to a maximum value of 1.65. This can be explained by an increasing number of sixfold rings in the sample and the onset of layer stacking, as confirmed by XRD and TEM results. A further increase in annealing temperature to 2000 C does not significantly change the number of sixfold rings, but leads to continued straightening and stacking of layers and to the formation of graphitic nanoribbons and nanocrystals, as shown in Fig. 1c.

3.3.

Porosity

The porosity of the vacuum-annealed TiC-CDC was studied using argon adsorption. Argon sorption isotherms (Fig 5a) show a complex dependence on annealing temperature. All are Type I of the Brauner classification [30], with the exception of the two samples annealed at 1800 and 2000 C. These two samples are markedly different because of large-scale structural collapses, which eliminated most pore volume. This behavior correlates well to effects previously observed for other non-graphitizable carbons annealed at temperatures that typically lead to graphitization [31]. The type I isotherms indicate a microporous material with substantial pore volume [1]. The total pore volume increases with increasing annealing temperature up to 1500 C, before

falling to very low values at 2000 C (Fig. 5a) as a result of structural collapse. The semi-log plot of the isotherms (Fig. 5b) shows the onset of adsorption beginning at lower relative pressures for samples annealed at lower temperatures. Samples with smaller pores adsorb larger quantities of gas at lower partial pressures [17]. The average pore size increases with increasing annealing temperature (Fig. 5c). The sample annealed at 2000 C did not have a sufficient porosity to calculate an average pore size from the isotherm. In general, the pore size changes from 0.66 nm for the assynthesized carbon to 1 nm for the CDC annealed at 1600 C. At higher temperatures (1800 C), increasing graphitization leads to pore coalescence and more than doubles the average pore size, reaching a value of 2.1 nm. The BET SSA, which is initially 1425 m2/g for the as-synthesized TiC-CDC, follows the same trend as the total pore volume (Fig. 5(c)). At annealing temperatures up to 1500 C, there is a fairly linear increase in SSA up to 2085 m2/g. At annealing temperatures above 1600 C, the SSA falls to a low value indicating that little porosity is accessible to the adsorbate molecules. This control over pore size and surface area via vacuum annealing is of particular interest, as it offers an alternative to the changes in CDC synthesis temperature [2] or activation [14]. Fig. 6a shows the SSA and pore size development of the same TiC-CDC upon isothermal activation at 875 C in CO2. While similar SSA values can be achieved via the oxidation in air, it leads to significant sample loss (>40%) and a substantial broadening of the pore size distribution, which is unfavorable for most CDC applications requiring well defined pores. Fig. 6b shows changes in SSA and pore size of TiC-CDC synthesizes at 1000 C during activation in air. It should be noted that higher synthesis temperatures yield CDCs with a higher degree of graphitization and larger average pore size, as compared to TiC-CDC produced at 600 C [2]. However, post-annealing of TiC-CDC-600 C at 1000 C under vacuum allows for similar SSA, but without a notable increase in pore size. Furthermore, activation of TiC-CDC-1000 C leads to a significant increase in SSA, but suffers from samples loss in increases in pore size, thus limiting microstructure control in CDC synthesized at higher chlorination temperature. Thus, by optimizing the annealing conditions, vacuum treatment can be used to produce highly porous carbon that exhibits both high SSA and narrow pore-size distributions without sample loss and surface functionalization. Vacuum

Fig. 5 – (a) Full-scale argon adsorption isotherm collected at 77 K. (b) Semi-logarithmic plot that shows low-relative pressure adsorption behavior. (c) BET SSA as a function of annealing temperature and average pore size calculated using NLDFT as a function of annealing temperature.

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Fig. 6 – Surface area and pore size development of CDC during gas phase activation. (a) Isothermal activation of TiC-CDC (synthesized at 600 C) in CO2 at 875 C. (b) Activation of TiC-CDC (synthesized at 1200 C) in air for 6 h at 415, 430, 450, and 475 C. The values reported in the brackets denote the weight loss that occurred during annealing. The ‘‘X’’ markings on y-axes represent the corresponding values for the as-produced TiC-CDC.

annealed CDCs may find applications in supercapacitor electrodes, gas chromatography, sorption, sensing and other applications when high SSA values are required and water adsorption in pores is undesirable. While this article focused on the carbon structure, properties of vacuum annealed CDC will be reported in a follow-up study, particularly its electrochemical characteristics and suitability for energy storage application, such as supercapacitor and Li-ion battery electrodes.

4.

Conclusions

Vacuum annealing of TiC-CDC leads to an increase in the pore volume and specific surface area with no significant change in the pore size up to 1500 C. This treatment allows synthesis of porous carbons with subnanometer pore size and SSA up to 2000 m2/g. Graphitization and collapse of the pore structure occur above 1500 C. Vacuum annealing of CDC therefore provides a suitable alternative to other structural modification techniques, such as activation, which suffer from high sample loss and extensive surface functionalization.

Acknowledgements The authors are grateful to P. Valenzuela and Dr. G. Yushin (currently at Georgia Tech) for experimental help with TEM analysis. We would also like to thank Dr. Ranjan K. Dash (YCarbon), Boris Dyatkin (Drexel University), and Dr. Patrice Simon (Universite´ Paul Sabatier, Toulouse, France) for helpful discussions. Research at Drexel University was supported by the U.S. Department of Energy, Office of Basic Energy Sciences, Division of Materials Sciences and Engineering under Award No. DE-FG02-07ER46473. R E F E R E N C E S

[1] Marsh H, Reinoso FR. Activated Carbon. Amsterdam: Elsevier; 2005. [2] Presser V, Heon M, Gogotsi Y. Carbide-derived carbons – from porous networks to nanotubes and graphene. Advanced Functional Materials 2011;21(5):810–33.

[3] Porada S, Weinstein L, Dash R, van der Wal A, Bryjak M, Gogotsi Y, et al. Water desalination using capacitive deionization with microporous carbon electrodes. ACS Applied Materials and Interfaces 2012;4:1194–9. [4] Gogotsi Y, Portet C, Osswald S, Simmons JM, Yildirim T, Laudisio G, et al. Importance of pore size in high pressure hydrogen storage by porous carbons. International Journal of Hydrogen Energy 2009;34(15):6314–9. [5] Zuckerbrod D, Kobayashi K. How Dry am I: Optimizing Cell Performance Through Proper Component Drying, 15th International Seminar On Double-Layer Capacitors and Hybrid Energy Storage Devices, 2005: Redox Engineering, Deerfield Beach, Florida. [6] Kinoshita K. Carbon, Electrochemical and Physicochemical Properties. New York: Wiley-Interscience; 1998. [7] Oschatz M, Kockrick E, Rose M, Borchardt L, Klein N, Senkovska I, et al. A cubic ordered, mesoporous carbidederived carbon, for gas and energy storage applications. Carbon 2010;48(14):3987–92. [8] Schmirler M, Knorr T, Fey T, Lynen A, Greil P, Etzold BJM. Fast production of monolithic carbide-derived carbons with secondary porosity produced by chlorination of carbides containing a free metal phase. Carbon 2011;49(13):4359–67. [9] Hossain MS, Tryk D, Yeager E. The electrochemistry of graphite and modified graphite surfaces – the reduction of O2. Electrochimica Acta 1989;34(12):1733–7. [12] Dash R, Chmiola J, Yushin G, Gogotsi Y, Laudisio G, Singer J, et al. Titanium carbide derived nanoporous carbon for energy-related applications. Carbon 2006;44(12):2489–97. [13] Kormann M, Popovska N. Processing of carbide-derived carbons with enhanced porosity by activation with carbon dioxide. Microporous and Mesoporous Materials 2010;130(1– 3):167–73. [14] Osswald S, Portet C, Gogotsi Y, Laudisio G, Singer JP, Fischer JE, et al. Porosity control in nanoporous carbide-derived carbon by oxidation in air and carbon dioxide. Journal of Solid State Chemistry 2009;182(7):1733–41. [15] Schmirler M, Glenk F, Etzold BJM. In situ thermal activation of carbide-derived carbon. Carbon 2011;49(11):3679–86. [16] Portet C, Kazachkin D, Osswald S, Gogotsi Y, Borguet E. Impact of synthesis conditions on surface chemistry and structure of carbide-derived carbons. Thermochimica Acta 2010;497(1–2):137–42. [17] Ravikovitch PI, Neimark AV. Characterization of nanoporous materials from adsorption and desorption isotherms. Colloids and Surfaces A-Physicochemical and Engineering Aspects 2001;187–188:11–21.

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CARBON

5 0 ( 2 0 1 2 ) 4 8 8 0 –4 8 8 6

[18] Aladekomo JB, Bragg RH. Structural transformations induced in graphite by grinding: analysis of 0 0 2 X-ray diffraction line profiles. Carbon 1990;28(6):897–906. [19] Obraztsova ED, Fujii M, Hayashi S, Kuznetsov VL, Butenko YV, Chuvilin AL. Raman identification of onion-like carbon. Carbon 1998;36(5-6):821–6. [20] Ferrari AC, Robertson J. Raman spectroscopy of amorphous, nanostructured, diamond-like carbon, and nanodiamond. Philosophical Transactions of the Royal Society of London Series A 2004;362(1824):2477–2512. [21] Ferrari AC, Robertson J. Resonant Raman spectroscopy of disordered, amorphous, and diamond like carbon. Physical Review B 2001;64(7):075414. [22] Shimodaira N, Masui A. Raman spectroscopic investigations of activated carbon materials. Journal of Applied Physics 2002;92(2):902–9. [23] Ferrari AC, Robertson J. Interpretation of Raman spectra of disordered and amorphous carbon. Physical Review B 2000;61(20):14095–107. [24] Thomsen C, Reich S. Double resonant raman scattering in graphite. Physical Review Letters 2000;85(24):15215–7.

[25] Dresselhaus MS, Dresselhaus G. Light Scattering in Solids III. Berlin: Springer-Verlag; 1982. [26] Solin SA, Caswell N. Raman-scattering from alkali graphiteintercalation compounds. J. Raman Spectrosc. 1981;10:129–35. [27] Urbonaite S, Juarez-Galan JM, Leis J, Rodriguez-Reinoso F, Svensson G. Porosity development along the synthesis of carbons from metal carbides. Microporous and Mesoporous Materials 2008;113(1–3):14–21. [28] Palmer JC, Llobet A, Yeon SH, Fischer JE, Shi Y, Gogotsi Y, et al. Modeling the structural evolution of carbide-derived carbons using quenched molecular dynamics. Carbon 2010;48(4):1116–23. [29] Mapelli C, Castiglioni C, Zerbi G, Mu¨llen K. Common force field for graphite and polycyclic aromatic hydrocarbons. Physical Review B 1999;60(18):12710–25. [30] Brunauer S. The Adsorption of Gases and Vapors. Princeton, New Jersey: Princeton University Press; 1945. [31] Pierson HO. Handbook of Carbon, Graphite, Diamond, and Fullerenes: Properties, Processing, and Applications. Park Ridge NJ: Noyes Publications; 1993.