Heat Treatment-Induced Structural Changes in SiC-Derived

From the PSD analysis presented later in this section, it appears that SiC-CDC ..... in TiC-CDC at 77 K. The results are reported in the Supporting In...
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J. Phys. Chem. C 2010, 114, 16562–16575

Heat Treatment-Induced Structural Changes in SiC-Derived Carbons and their Impact on Gas Storage Potential Mauricio Rinco´n Bonilla, Jun-Seok Bae, T. X. Nguyen, and Suresh K. Bhatia* School of Chemical Engineering, The UniVersity of Queensland, St. Lucia, Brisbane QLD 4072, Australia ReceiVed: June 14, 2010; ReVised Manuscript ReceiVed: August 23, 2010

We investigate the effect of heat treatment on the structure of carbide-derived carbons (CDC) prepared by chlorination from nanosized βSiC particles and on their methane as well as hydrogen storage and delivery performance. Pore size and pore wall thickness distributions of the CDCs are obtained from interpretation of argon adsorption data using the finite wall thickness (FWT) model. The adequacy of the FWT model for adsorption modeling in the SiC-CDC samples is demonstrated by satisfactory prediction of subatmospheric and high pressure adsorption isotherms of CO2 and CH4 at 313 and 333 K. From the characterization results, it is observed that the SiC-CDC particles are predominantly amorphous with slight graphitization of the external surface. The degree of graphitization is more pronounced in the sample prepared at 1000 °C and increases slowly with heat treatment time. During this time the accessibility of methane molecules is found to increase, as a result of short-range ordering and opening up of pore entrances. Nevertheless, methane storage capacity is unsatisfactory, despite the high surface area and porosity, due to accessibility problems. On the other hand improvement in high pressure H2 uptake (4.61 wt % at 77 K) is obtained for SiC-CDC chlorinated at 800 °C and heat treated for one day. The recently predicted optimal delivery temperature of 115 K for hydrogen storage is found to be appropriate for this material. It is demonstrated that accessibility is an important issue to be addressed for methane storage in carbons, but which has hitherto not received attention for this application. 1. Introduction The energy and environmental problems arising from the use of conventional energy sources like petroleum and coal has led to considerable interest in hydrogen and methane. Hydrogen possesses a high chemical energy and pollution free burning. However, hydrogen manufacture and storage continue to be major issues for its adoption as a viable green fuel. On the other hand, methane is highly abundant as natural gas and biogas; its emissions are relatively low and it is the most likely candidate to replace the liquid fuels in the years to come. Nevertheless, the current mobile storage technologies for methane, like liquefaction or compression of natural gas, are either expensive or inefficient in terms of energy density. To solve the open problem of efficient storage of methane and hydrogen, adsorption in porous solids has gained attention as an attractive alternative due to the higher density of the adsorbed gas compared to that of bulk conditions. In recent years, there has been intensive research on the development of new materials for energy storage applications with the appropriate balance between cost and storage capacity.1 Among the options, carbidederived carbons (CDCs) have come up as an appealing choice due to their high surface area, well tunable pore size and ease of synthesis from natural carbides.2-4 One of the most interesting features of CDCs is that their pore structure, and thereby their properties, can be relatively well controlled by selecting the appropriate precursor, synthesis and post treatment conditions. To tailor carbons with optimal characteristics for maximum uptake of desired adsorbates (i.e., methane or hydrogen), considerable effort has been devoted to the understanding of the microstructural changes occurring during the preparation and post-treatment of CDCs from a number of precursors (e.g., * To whom correspondence may be addressed. E-mail: s.bhatia@ uq.edu.au.

SiC, TiC, VC, ZrC).3,5-9 In this task, a combination of techniques such as X-ray diffraction (XRD), transmission electron microscopy (TEM), Raman spectra, electron energy-loss spectroscopy (EELS), and adsorption using Ar and N2 at subcritical conditions have been employed.10-13 Some work on atomistic reconstruction using reverse Monte Carlo simulation of TiC-based carbon has been reported.14 Among these techniques, the adsorption method has played an important role in the search for the optimal structure for largest energy storage capacity. Nevertheless, there has been little work on the influence of preparation conditions and heat treatment on pore network accessibility and its effects in energy storage capacity of carbons. In this work, we show that a combination of the above analysis techniques and adsorption modeling can provide good insight into the microstructural changes during heat treatment of CDCs, and their influence on the storage capacity of methane and hydrogen. In previous work, the comparison of experimental adsorption data and modeling results for subatmospheric and supercritical adsorption of CO2 and CH4 in Ti3SiC26 and SiC11 derived carbons chlorinated at several temperatures, along with the information obtained from transmission electron microscopy permitted the extraction of qualitative information on the variation of pore accessibility (e.g., elimination of pore constrictions, graphitic pore wall thickening) and its relation with the dimensions and geometry of the adsorbate molecules. Pore network accessibility was shown to have a significant impact in adsorption behavior and, therefore, a major role is expected in the storage capacity of CDCs. Here, we explore the microstructural rearrangements occurring during heat treatment of SiC-CDCs prepared in our laboratory, and their influence on the storage of methane and hydrogen. The characterization was performed using helium

10.1021/jp105473x  2010 American Chemical Society Published on Web 09/13/2010

Structural Changes in SiC-Derived Carbons pycnometry, XRD, Raman spectroscopy, and high-resolution TEM (HRTEM). In addition, the finite wall thickness (FWT) model,15,16 utilizingnonlocaldensityfunctionaltheory(NLDFT),17,18 is used for interpretation of Ar adsorption at 87 K and for the prediction of experimental adsorption isotherms of CO2 and CH4 in the SiC-CDC. The FWT model was previously shown to be a suitable model for the characterization of activated carbons16 and for virgin SiC-CDC.11 Therefore, its performance is expected to be appropriate for heat-treated forms of SiC-CDC. Significant overprediction is observed in adsorption modeling of CH4, indicating that part of the pore volume accessible to the probe gas is inaccessible to methane. This has important consequences on the storage capacity of the SiC-CDCs, and we show that not only surface area and porosity are important issues to address when new materials are to be evaluated for methane storage, but the role of pore accessibility is also crucial. In theory, a H2 uptake of 4.61 wt % at 77 K and 28 bar is obtained from SiC-CDC chlorinated at 800 °C and heat treated for one day, which is very competitive compared to state of the art activated carbons19 and CDCs.5,7 While the theoretical predictions obtained here must be further corroborated by experiment, they provide good insight on the possibilities of such porous materials for H2 storage before performing intensive experimental work in the laboratory. The relationship between the structural changes, analyzed by mean of the characterization results, and the H2 storage capacity of the material is highlighted in Section 4.3.1. It is seen that the concepts of short and long ranged C-C restructuring, useful to explain the results of experimental analysis techniques, are also crucial to understand the behavior of the storage capacity with heat treatment. Finally, a discussion of the predicted solid deformation during high pressure adsorption of methane and carbon dioxide is presented. 2. Experimental Section The SiC-CDCs samples used in this study were synthesized in our laboratory by chlorination of commercial nanosized SiC powders (Sigma-Aldrich) at 800 and 1000 °C. The chlorination was performed on the SiC powders using ultrahigh purity chlorine (BOC Gases, 99.9%) as the reactive gas and ultrahigh purity argon (BOC Gases, 99.999%) as the purging gas. For the CDC synthesis, the preparation procedure is detailed elsewhere.20 For synthesis, the SiC powder was placed in a quartz sample boat and loaded into a horizontal quartz tube furnace. Initially, the tube with the quartz sample boat inside was purged with argon. Subsequently, the temperature of the furnace was raised until the desired chlorination temperature was reached. Argon purge was maintained through the tube furnace during the heating process, and once the desired temperature was attained and stabilized, the argon purge was stopped and replaced with pure chlorine flow to begin the chlorination process. During the chlorination, pure Cl2 gas flowed over the SiC sample at the rate of 50 mL per minute until completion of the chlorination reaction. When the chlorination was complete, Ar flow was used to cool the sample down to ambient temperature as well as to flush metal chlorides remaining in the sample. The residual chlorine and metal chlorides were captured in sodium hydroxide solution. Finally, the resulting CDC was heat treated for 1 and 3 days at 1100 °C in a pure argon atmosphere with a constant flow of pure argon of 50 mL per minute. For simplicity, the short name “SiC-CDC X HY” will be used for SiC derived carbons throughout the text. Here, X denotes the synthesis temperature of the CDC, while Y depicts heat treatment time. For instance, SiC-CDC 1000 H1 represents a CDC sample prepared at a synthesis temperature

J. Phys. Chem. C, Vol. 114, No. 39, 2010 16563 of 1000 °C and followed by a 1 day heat treatment at 1100 °C. The same notation rule applies for the heat-treated forms of Ti3SiC2 mentioned throughout this work for comparison purposes (i.e., Ti3SiC2-DC 1000 H1). For the characterization of the CDCs, the skeletal density of the SiC-CDC samples was measured by means of buoyancy analysis under He at a temperature of 383 K over a pressure range of 4 to 25 bar,21 using our gravimetric sorption system (Rubotherm Pra¨zisionsmesstechnik GmbH, Bochum, Germany). These measurements yielded a linear relation between the change in solid mass and He bulk density, from whose slope the skeletal density of the CD was determined, as discussed elsewhere.21 Measurements were also performed with helium pycnometer (Micromeritics Accupyc 1330), obtaining a difference within 2%. The results reported here are those obtain from the former method. XRD analysis was performed using Cu KR radiation (40 kV, 40 mA, l ) 1.54056 Å) with a step size of 0.02° (2θ) and fixed slit mode over the range 10-95°. Raman spectroscopy was measured using a Renishaw Ramanscope Raman spectrometer, utilizing 1 K series He-Ne laser excitation (632.8 nm). The local two-dimensional atomistic structure of SiC-CDC samples was analyzed from HRTEM images obtained using a TECHNAI F30GTEM cryo-electron transmission electron microscope with optimal resolution of 1.5 nm at 300 kV. The pore structure analysis of the SiC-CDC samples was performed by interpretation of the argon adsorption isotherm at 87 K using the FWT model.15,16 For adsorption measurements, a Micromeritics ASAP2020 volumetric adsorption analyzer was used to obtain the adsorption isotherm data of argon at 87 K and CO2 at 273 K. High-pressure adsorption data of CO2 and CH4 at 313 and 333 K up to 200 bar was obtained using a gravimetric sorption system (Rubotherm Pra¨zisionsmesstechnik GmbH, Bochum, Germany). The detailed description of the high-pressure adsorption measurement procedure is given elsewhere.11,22 3. Mathematical Model The virgin and heat-treated forms of SiC-CDC have been characterized for the pore size distribution (PSD) and pore wall thickness distribution (PWTD) by interpretation of Ar adsorption data at 87 K using the Finite Wall Thickness model recently developed in our laboratory.15,16 With these characterization results, the experimental adsorption isotherms of CO2 and CH4 were predicted. The theoretical excess amount adsorbed was calculated assuming a flexible carbon structure following a recent work of this group.23 In summary, it is given by

mfex ) mfa -

mHe a FHe b

Ffb - ∆VmaxFfb

(1)

f is the theoretical excess amount adsorbed, Fbf is where mex He the adsorbate bulk density, mHe is the pore volume a /F b accessible to helium and mfa is the absolute adsorbed amount calculated by the FWT model assuming a rigid pore structure, following

mfa )

∫ Fˆ (P, Hin)f(Hin)dHin

(2)

with f(Hin) being the PSD of the adsorbent on the geometrical pore width Hin ) Hcc - σc, where Hcc is the center-to-center pore width and σc is the effective carbon diameter. Fˆ (P,Hin)

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is the average adsorbate density at pressure P in a pore of width Hin and is given by ∞

Fˆ (P, Hin) )

The interaction potential φsf(r) between a molecule and a carbon wall with n graphene layers is represented by28

∞

n-1

∑ p(m) ∑ p(l) H1in ∫ Flm(P, Hin, z)dz

m)1

φsf(r) ) 2πFscσsf2εsf

l)1

i)0

(3) where p(n) denotes the probability of having a wall of n graphene layers, and Flm(P,Hin,z) is the local density profile calculated with the Tarazona prescription of the nonlocal density functional theory17,18 in a pore with m and l graphene layers in the left and right walls, respectively. As shown in He previous work,23 the term mHe a /F b in eq 1 can be replaced by the pore volume obtained from interpretation of argon adsorption data at 87 K using the FWT model. Accordingly, He for this paper the term mHe a /F b will be replaced by the pore volume estimated with the FWT from argon adsorption at 87 K. Finally, the term ∆Vmax in eq 1 denotes the maximum difference between the external solid volume in helium and analysis gas atmosphere. For a given porous carbon, ∆Vmax is clearly a function of the adsorbate accessibility. At high pressures, some of the pore bodies accessible to helium but inaccessible to the adsorbate molecules can be compressed, leading to shrinkage of the carbon particles under the action of the external bulk pressure. On the other hand, strong adsorption can lead to swelling of the adsorbent. In practice, the deformation ∆Vmax can only be calculated by fitting the theoretical adsorbed amount to the experimental data measured over a wide range of high pressure values at which the bulk density approaches the adsorbed phase density. Equation 1 has two important advantages. (i) It avoids the uncertainty arising from the location of the Gibbs dividing surface between the adsorbent and the adsorbed phase, which is especially hard to define in materials with a significant nanosized porosity, and (ii) it is not sensitive to the location of the boundary separating the bulk phase and the external adsorbent volume, since the external surface area is negligible compared to the internal surface area. For the prediction of the single pore density Fˆ (P,Hin), the interaction between fluid molecules was represented by the classical Lennard-Jones (LJ) 12-6 potential

[( ) ( ) ]

φff(r) ) 4ε

σff r

12

-

σff r

6

(4)

where r is the interparticle distance, σff is collision diameter and εff is well depth. For the particular case of H2, where quantum effects are important, the semiclassical Feynman-Hibbs (FH)24 potential was employed

φHff 2(r) ) φff(r) +

( )(

βp2 2 ∂φff(r) ∂2φff(r) + 12mH2 r ∂r ∂r2

)

∑

(5)

where mH2 is the hydrogen molecule mass. Higher order corrections are neglected since they are insignificant at temperatures of 77 K or above, relevant to this work.25 For the calculation of the excess free energy functional17,18 of H2, the Weeks, Chandler, and Andersen (WCA)26 approach was used along with the scheme suggested by Kowalczyk et al.27 to describe the attractive part of the FH potential.

[(

2 σsf 5 z + i∆

) ( 10

-

σsf z + i∆

)] 4

(6) Here z is the center to center distance between the fluid molecule and the pore wall surface, ∆ is interlayer spacing, and Fsc is the number of carbon atoms per unit area in a single graphene layer. Quantum effects must be accounted for in the case of H2, where the FH correction yields

φHsf2(z) ) φsf(z) +

( )

βp2 ∂2φsf(z) 24µ ∂z2

(7)

The external potential profile, φext(z) for a slit-shaped pore having a center to center size H is determined from superposition of the potentials of the opposing pore walls

φext(z) ) φsf(z) + φsf(H - z)

(8)

φHext2 (z) ) φHsf2(z) + φHsf2(H - z)

(9)

Thus, for H2

Finally, the hydrogen density profile resulting from the solution of the Euler-Lagrange equations arising from the classical NLDFT, must be convoluted with the three-dimensional wave-packet.25 The combination of NLDFT and the FH corrections will be called FH-NLDFT. 4. Results and Discussion 4.1. Characterization. In this section, we present the characterization results of the virgin and heat-treated forms of the SiC-CDC prepared from 50 nm particles by chlorination of βSiC particles at 800 and 1000 °C, according to the procedure detailed in Section 2. The structural changes occurring during the samples heat treatment are due to short ranged C-C interactions (chemical) and long ranged interactions (physical). To clarify the structural effects arising from these two types of interaction, we briefly reproduce their definitions from our recent work.6 “Short ranged C-C restructuring” arises from chemical interaction between neighboring atoms, which tends to reduce active sites or defects, leading to only a subtle change in the size of pore bodies but significant modification of pore network connectivity. On the other hand, “long-ranged restructuring” arises as a result of van der Waals interaction between two or more nearest neighboring carbon walls, leading to the formation of thicker walls (i.e., wall coalescence), resulting in the population of small pores (6 Å) increases. These two concepts will be used in the subsequent sections for interpreting the effects of heat treatment on the microstructural evolution of SiC-CDCs. It is important to remark that wall coalescence does not necessarily mean the collapse of graphene sheets, but the approaching of opposing pore walls to an extent which, as a consequence of wall irregularities, prevents the diffusion of gas molecules in the space between the walls.

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Figure 1. Helium skeletal density of virgin and heat-treated forms of SiC-CDC (circles) and Ti3SiC2-CDC (triangles) chlorinated at 800 °C (dashed line) and 1000 °C (solid line).

4.1.1. Helium Density. Helium skeletal density is depicted in Figure 1 for SiC-CDC 800, SiC-CDC 1000 and their heattreated forms. For comparison purposes, the helium skeletal density of virgin and heat-treated forms of Ti3SiC2-derived carbon taken from our previous work6 are also shown in Figure 1. For both the virgin SiC-CDC 800 and the SiC-CDC 1000, the initial density is significantly higher than the graphite density (∼2.27 g/cm3), indicating the presence of an important population of sp3 C-C bonds after chlorination. It is interesting to see that the density of the virgin carbons synthesized at 800 and 1000 °C are very similar (∼2.5 g/cm3), suggesting that the sp3/sp2 ratio is not considerably affected by the chlorination temperature, but depends mostly on the sp3/sp2 ratio of the precursor SiC. This idea is supported by Urbonite et al.,12 who used EELS for the determination of the sp3/sp2 ratio of TiCderived carbons synthesized at several temperatures. Nevertheless, the helium density of the virgin Ti3SiC2-CDC 800 and Ti3SiC2-CDC 1000 are slightly different (2.2 and 2.3 g/cm3, respectively), which apparently suggests that the sp3/sp2 ratio increases with the chlorination temperature for this material. However, this is not associated with a higher degree of sp3hybridization at higher chlorination temperatures, but with the enhancement of the accessible pore volume for He.6 The fact that He accessibility is affected by the chlorination temperature in Ti3SiC2-CDC but remains unchanged for the SiC-CDC can be possibly due to the difference in the microstructural rigidity of both materials. After the Si and Ti atoms are leached from the precursor Ti3SiC2, a layered carbon structure of predominantly sp2 hybridized carbon atoms remains.13 As chlorination takes place, the regions in the material that have been already leached may undergo significant structural changes, especially in the directions perpendicular to the carbon layers. These structural rearrangements are encouraged by higher chlorination temperatures, leading to variations in He accessibility for the virgin samples and thus to differences in the skeletal density. On the other hand, the βSiC consists of a tetrahedral network of silicon and carbon atoms with sp3 hybridization.4 After leaching, there is no preferential bonding direction and a highly isotropic structure results.3,11 Since this structure has less degrees of freedom than Ti3SiC2-CDC its rigidity is higher, and the structural rearrangements occurring during chlorination are not significant enough to cause significant variations in He accessibility for the virgin SiC-CDC chlorinated at different temperatures.

Figure 2. XRD pattern of the virgin and heat-treated forms of SiCCDC chlorinated at (a) 800 and (b) 1000 °C. The increasing intensity of the 100 reflection indicates higher 2D-regularity.

For both the SiC-CDC and the Ti3SiC2-CDC, the helium density increases with heat treatment time. Surprisingly, the increase in helium density is significantly more pronounced for the SiC-CDC 800 than it is for the SiC-CDC 1000 after three days of heat treatment. This contrasts with the trend observed in the Ti3SiC2-CDC samples, in which heat treatment leads to an increase in He density at a similar rate for both the samples synthesized at 800 and 1000 °C. Ti3SiC2-CDC undergo extensive graphitization after heat treatment.6 However, the SiC-CDC 1000 and SiC-CDC 800 apparently show different evolution paths. From the PSD analysis presented later in this section, it appears that SiC-CDC 1000 structure is more rigid and tends to present the wall coalescence phenomenon from the early stages of heat treatment, whereas for SiC-CDC 800 the short-range rearrangements dominates during a long part of the heat treatment process, leading to a significant increase in the pore volume accessible to He. 4.1.2. X-ray Diffraction. The XRD pattern of the virgin and heat-treated forms of SiC-CDC 800 and SiC-CDC 1000 are depicted in Figure 2. The 002 reflection is very broad in all the samples, revealing poor degree of graphitization. On the other hand, SiC-CDC 800 shows a clear increase in the intensity of the 100 reflection after one and three days of heat treatment, suggesting increase in the two-dimensional regularity29 of the structure. Since the 002 reflection reveals no parallel stacking of graphite layers, it is likely that randomly distributed pieces of graphitized carbon sheets are becoming larger and more symmetric. As discussed below, the TEM micrographs showed these graphitic sheets to be predominantly located in the shell

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Figure 3. TEM micrographs of the virgin and heat-treated forms of SiC-CDC synthesized at 800 and 1000 °C. The circles depict areas in the particles shell where some graphitization occurs. The core of the carbons remains predominantly amorphous.

region. The 100 reflection is more spread out in the SiC-CDC 1000 and its heat-treated forms, although an increase in the intensity of this peak with heat treatment duration is also evident. 4.1.3. TEM. The TEM images of the SiC-CDC 800, SiCCDC 1000, and their heat-treated forms are depicted in Figure 3. These images reveal a predominantly amorphous structure with formation of a graphitic shell on the SiC-CDC particles as the heat treatment is carried out. The formation of a shell of turbostratic carbon in SiC-CDC was reported by Welz et al.3 for chlorination temperatures below 1000 °C and etching with pure chlorine. This is probably due to the break of the symmetry around the carbon atoms located in the particles surface, which are not free to form sp3 bonds in all directions of the space.11 It can be observed that the graphitization in

the outer layer of the carbon particles is favored by heat treatment, but no long-range parallel layering occurs. This is in agreement with the results obtained from the XRD pattern, suggesting two-dimensional growth of graphitic sheets but no significant stacking. The absence of long-range parallel stacking is possibly due to the presence of an intricate interface with rigid sp3 hybridized carbon segments, which do not permit the graphitic sheets to be accommodated in their equilibrium parallel configuration. 4.1.4. Raman Spectra. Figure 4 depicts the Raman Spectra of SiC-CDC 800, SiC-CDC 1000 and their heat-treated forms. The ratio of the “disorder band” (D) to the “graphite band” (G) is plotted in the insets of Figure 4 panels a and b. According to Ferrari and Robertson,30 an increase of the I(D)/I(G) ratio for

Structural Changes in SiC-Derived Carbons

Figure 4. Raman spectra of the virgin and heat-treated forms of SiCCDC chlorinated at (a) 800 and (b) 1000 °C. The D and G indicate the position of the disordered and graphite bands respectively. Inset: variation of I(D)/I(G) ratio with heat treatment time.

an amorphous carbon is an indication of ordering when the graphite sheets have an average diameter La below 20 Å. Above La ) 20 Å, the ordering is accompanied by a decrease in the I(D)/I(G) ratio with increasing La. For the SiC-CDC 800, an increasing I(D)/I(G) is observed. Since the TEM and XRD reveal ongoing graphitization in the surface and consequently, structural ordering, the size of the sheets should be below 20 Å. In the case of SiC-CDC 1000 the I(D)/I(G) ratio goes through a maximum after one day heat treatment, indicating probably a transition between an average diameter below 20 Å to values above 20 Å. The small La extracted from the model of Ferrari and Robertson confirms that these structures are highly amorphous with graphitization limited to local areas in the structure. 4.1.5. Pore Size Distribution. The pore size distribution of the SiC-CDC 800 and its heat-treated forms obtained with the FWT model is depicted in Figure 5a. These PSD results were determined by interpretation of Ar adsorption experimental data at 87 K. In Table 1, the surface area, pore volume and average pore size obtained from the FWT model are reported. After one day heat treatment, all the major peaks at Hin > 5 Å become more intense, but a slight decrease in the intensity of the first peak (Hin < 0.5 Å) is also observed. However, the cumulative distribution shown in the inset of Figure 5a reveals that the broadening of the first peak width leads to a net increment of the pores population below 5 Å. These structural changes are summarized in Table 1, where a very significant increase in the volume accessible to Ar at 87 K, along with an important enhancement of the surface area is observed after one day heat

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Figure 5. PSD of the virgin and heat-treated forms of SiC-CDC chlorinated at (a) 800 and (b) 1000 °C. The PSD is obtained by interpretation of experimental argon adsorption isotherms at 87 K by means of the FWT model. The inset in (a) depicts the cumulative pore volume distribution of SiC-CDC chlorinated at 800 °C, and of its heattreated forms.

TABLE 1: Surface Area, Pore Volume, and Average Pore Width of the Virgin and Heat-Treated SiC-CDCs Obtained from the FWT Model from Interpretation of Ar Adsorption at 87 K sample SiC-CDC SiC-CDC SiC-CDC SiC-CDC SiC-CDC

800 800 H1 800 H3 1000 1000 H3

surface area (m2/g)

pore volume (cm3/g)

average pore widtha (Å)

1929.3 2301.5 2139.7 1927.6 1908.3

0.601 1.039 1.064 0.639 0.908

7.15 8.17 8.55 7.85 8.05

a The average pore size is defined as ∫Hf(H)dH/Vp, where Vp is the sample porosity.

treatment. These changes in the pore structure can be understood in terms of the competing effect of the short-range C-C restructuring that tends to increase pore network accessibility, and the long-range interactions due to van der Walls forces that lead to coalescence of small pore bodies (Hin < 5 Å). After one day heat treatment, the C-C restructuring predominates, causing a net increase of both surface area and accessible pore volume for the probe gas. Long-range interactions lead to walls collapse in pores with size below 4.6 Å, as it is observed from the cumulative distribution in the inset of Figure 5a. However, its impact is overcome by the overall increase in the accessibility

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arising from the short-range rearrangements. It is interesting to see that a shoulder appears in the PSD of the SiC-CDC 800 H1 below 4.2 Å. This is an indication of the highly amorphous nature of the structure.31 As the pore walls come closer due to van der Waals forces, the walls irregularities create a spectrum of narrow spots within the pore bodies that is reflected in the PSD by the appearance of the shoulder in the very small width region. The PSD of the SiC-CDC 800 H3 depicted in Figure 5a shows a dramatic decrease in the intensity of the first peak below 5 Å compared to the SiC-CDC 800 H1 and the virgin carbon. Concurrently, all the peaks above 5 Å become more intense. Moreover, Table 1 reveals that between one and three days of heat treatment a very slight increase in the volume accessible to Ar at 87 K occurs, along with an important decrease in the surface area. This is a consequence of the wall coalescence phenomenon arising from long-range rearrangements. As the narrower pores disappear, their volume is inherited by the neighboring pores, leaving the total porosity almost unchanged but causing a great decrease in the surface area. Wall coalescence is encouraged by the short-range interactions taking place mainly at the beginning of the heat treatment. As local defects are eliminated, the 2D regularity within the pore walls is increased, as shown by the XRD pattern in Figure 2. The higher regularity enhances the impact of van der Waals forces between sufficiently close walls and, as a consequence, the walls collapse is promoted. Nevertheless, the structure still remains highly disordered, as is clear from the TEM images in Figure 3. This can be verified in the PSD by noting that the shoulder below 4.2 Å in the PSD of the SiC-CDC 800 H3 becomes more pronounced and reaches lower pore sizes compared to SiC-CDC 800 H1. The slightly higher regularity of the carbon sheets structure allows the neighboring walls to come closer under the effect of the van der Walls forces, creating a larger amount of narrower spots within the pore bodies that is reflected in the growth of the shoulder. For fluid molecules such as CH4 with larger effective diameter than argon, the structural variations reflected in the PSD have interesting implications on the accessibility, as will be further shown. Figure 5b depicts the PSD of SiC-CDC 1000 and its three days heat-treated form, SiC-CDC 1000 H3. Very subtle variations occur, although a slight decrease in the surface area accompanied by an increase in the pore volume can be seen in Table 1. The increase in porosity comes mainly from an increase in the intensity of pores above 8 Å, while a slight decrease in the intensity of pores below 8 Å leads to a reduction of surface area. This is an indication of wall coalescence, discussed above for the SiC-CDC 800 and heat-treated samples. However, the magnitude of the microstructural variations is less dramatic than the ones observed in the case of SiC-CDC 800, in agreement with the small changes registered in the He density for the heattreated forms of SiC-CDC 1000. The fact that the surface area of SiC-CDC 1000 H3 is lower than the surface area of the virgin SiC-CDC 1000, whereas the opposite trend is obtained for SiCCDC 800 and SiC-CDC 800 H3, supports the idea that van der Wall forces are the dominant effect driving the microstructural transformations during heat treatment of SiC-CDC 1000. A correlation between average pore size and H2 storage capacity has been shown to exist32 with pores between 6 and 8 Å having the highest uptake per unit surface area. Theoretical approaches locate the optimal pore size around 5-6 Å33 or 9.2 Å.34 In particular, the value 9.2 Å has been derived from purely geometrical considerations with no assumption on the form and parameters for the C-H2 interaction potential. In any case, it is

Bonilla et al.

Figure 6. Predicted and experimental subatmospheric adsorption isotherms of CO2 at 273 K on the SiC-CDC 800, SiC-CDC 800 H1, and SiC-DC 800 H3.

TABLE 2: Interaction Parameters Used within the NLDFT Scheme for the Prediction of the Adsorption Isotherms of the Methane, Carbon Dioxide, and Hydrogen in the SiC-CDC Samples gas

σff(Å)

εff/kB(K)

σcf(Å)

εcf/kB(K)

source

CH4 CO2 H2

3.6177 3.472 2.958

146.91 221.98 36.7

3.509 3.436 3.178

64.14 78.84 36.98

34 34 22, 39

clear that average pore sizes below 10 Å are convenient for H2 storage, and this condition is perfectly met by the SiC-CDC and its heat-treated forms. The theoretical results obtained here for H2 tend to support the idea of an optimum average pore size closer to 9.2 Å, as will be discussed later. 4.2. Adsorption Modeling of CO2. Figure 6 depicts the predicted and experimental adsorption isotherms of CO2 at 273 K and subatmospheric pressure in SiC-CDC 800 (s), SiC-CDC 800 H1 (---), and SiC-CDC 800 H3 (- · -). The prediction was made by means of the FWT model assuming a rigid carbon structure (∆Vmax ) 0). The excess theoretical adsorption isotherm was calculated using eq 1. For the NLDFT calculation of the local adsorption isotherm in a single pore, the LJ CO2-CO2 parameters used in a recent paper from our group31 and summarized in Table 2 where used. The LJ C-C interaction parameters are taken to be similar to those of graphite: σc ) 3.4 Å and εc/kBT ) 28 K. The fluid-carbon interaction parameters are determined from the Lorentz-Berthelot combining rule. The binary parameter kfc,31 introduced to account for packing effects, is taken to be zero at subatmospheric pressures, since the one center assumption for the CO2 molecules has been found to be appropriate at such low densities.11 Figure 6 shows that the prediction of the adsorption isotherm of CO2 in the virgin SiC-CDC 800, although not accurate, is very good considering the predominantly amorphous structure of the material. With longer heat treatment times, the model prediction becomes better, which is in agreement with the slow increase in structural regularity observed in the XRD pattern for the heat-treated SiC-CDC 800 samples, making the slit pore approximation with perfectly flat graphite walls increasingly precise. It is interesting to see that the predicted amount adsorbed decreases with the heat treatment duration at almost every pressure in the range, despite the large increase in the surface area and the porosity, particularly for the SiC-CDC 800 H1. The reason is that for the highly amorphous SiC-CDC 800 the microporous region accounts for most of the volume accessible

Structural Changes in SiC-Derived Carbons

Figure 7. High-pressure experimental and predicted adsorption isotherms of CO2 in (a) SiC-CDC 800, (b) SiC-CDC 800 H1, and (c) SiC-CDC 800 H3 at 313 and 333 K.

to the probe gas. In particular, pores under 6 Å constitute more than 50% of the measured porosity. In contrast, for SiC-CDC 800 H1 and SiC-CDC 800 H3 pores under 6 Å make up to 28 and 23% of the measured porosity. As it is well-known, at low pressure adsorption is governed by the solid-fluid interactions,35 which are highly enhanced in the very narrow pores due to the overlapping of the potential wells from opposite walls. Consequently, the reduction in the proportion of very narrow pores results in a decrease on the amount adsorbed at low pressure for the heat-treated forms of SiC-CDCs. Figure 7a-c depicts the predicted and experimental high pressure adsorption isotherms of CO2 in SiC-CDC 800 and its heat-treated forms at 313 K (s) and 333 K (---). The predicted values of the excess amount adsorbed were obtained from eq

J. Phys. Chem. C, Vol. 114, No. 39, 2010 16569 1, assuming a nonrigid (∆Vmax * 0) structure for the carbon samples. ∆Vmax represents the change in the external volume of carbon particles in the analysis gas atmosphere relative to that in a He atmosphere.23 In this regard, the nonzero value of ∆Vmax accounts for the deviation between the rigid carbon model based theoretical isotherm (∆Vmax ) 0) and the experimental data. ∆Vmax is determined by fitting the predicted excess amount adsorbed to the high pressure branch of the experimental isotherm. The binary interaction parameter was taken to be kfc )0.09367 for the NLDFT calculations, following prior work from this laboratory.31 This value for kfc has been shown to correct appropriately the effects of the one center approximation used here to model the CO2 molecule, as discussed elsewhere.11 Figure 7a shows that the predicted isotherm of CO2 in the highly amorphous virgin carbon, SiC-CDC 800, differs slightly from the experimental high-pressure adsorption data. This is not surprising, since the structure of the sample deviates strongly from the infinite slit-shaped pores assumption. However, it is readily seen from Figure 7b,c that the quality of the prediction improves with longer heat treatment times in analogy with the case of subatmospheric adsorption of CO2 at 273 K. This is most likely a consequence of the increase in regularity of the pore walls, as discussed before. It is important to remark that the good agreement between predictions and experiments indicates that the PSD obtained by interpretation of argon experimental adsorption at 87 K can appropriately capture the pore volume accessible to CO2 on each of the SiC-CDC samples. The accessibility of CO2 is comparable to the H2 accessibility, and the prediction of H2 uptake at 77 K can be successfully done by FH-Grand canonical Monte Carlo (GCMC) simulation using the CO2 derived PSD if the appropriate H2-C interaction parameters are chosen.22 Since the pore volume obtained by Ar 87 K characterization coincides with the volume accessible to CO2 at 298, 313, and 333 K, it is expected that the H2 uptake can be realistically predicted from the Ar 87 K characterization of the SiC-CDCs. This will be further discussed in Section 4.3.1 4.3. Energy Storage. 4.3.1. Hydrogen Storage in HeatTreated SiC-CDC. In this section, we investigate the suitability of heat-treated SiC-CDC for hydrogen storage purposes. The use of SiC-CDC for H2 storage had been previously explored using inactivated SiC-CDC7 as well as H2-annealed ordered mesoporous (OM) SiC-CDCs.36 However, to our knowledge, there is no reported information on the effect of heat treatment on the H2 storage capacity of CDCs. To assess the viability of heat-treated forms of SiC-CDC for H2 storage, the Tarazona NLDFT17,18 scheme along with the Feynman-Hibbs (FH) corrections (FH-NLDFT)24 detailed in Section 3 was used to predict the H2 adsorption in a single slit shaped pore. Equation 1 was then used to predict the excess amount adsorbed in the SiC-CDC samples taking ∆Vmax equal to zero, since H2 is weakly adsorbed and its molecular dimensions are comparable to those of He, suggesting that the volume accessible to both species is similar and the extra strain caused by H2 adsorption in the pressure range under study (