J. Phys. Chem. C 2008, 112, 17427–17435
17427
Role of Carbon Order in Structural Transformations and Hydrogen Evolution Induced by Reactive Ball Milling in Cyclohexene Apurba Sakti,† Nichole M. Wonderling,‡ Caroline E. B. Clifford,§ John V. Badding,| and Angela D. Lueking*,† Department of Energy & Mineral Engineering, Materials Research Institute, EMS Energy Institute, and Department of Chemistry, The PennsylVania State UniVersity, 120 Hosler, UniVersity Park, PennsylVania 16802 ReceiVed: June 30, 2008; ReVised Manuscript ReceiVed: August 18, 2008
Demineralized Summit (DS) anthracite, DS annealed at 1673 K, and graphite are used to explore the effect of precursor order on structural transformations and H2 evolution that result during reactive ball milling. Carbon structure was assessed before and after milling with temperature-programmed oxidation, X-ray diffraction (XRD), ultraviolet Raman spectroscopy, N2 adsorption, He density, and solvent swelling. Graphite milled in cyclohexene is primarily nanocrystalline graphite, with 8 wt % amorphous content leading to lowtemperature oxidation, swelling, increased surface area, and mesoporosity. Milling the disordered DS leads to signs of increased sp2 clustering, increased cross-linking, a significant ultramicroporosity with pores less than 8 Å, and low-temperature H2 evolution. The carbon fraction of annealed DS behaves similarly to graphite in the mill. Introduction Ball milling of graphite has been used as a strategy to produce carbon-based hydrogen storage materials.1-9 We recently reported that substitution of anthracite coal for graphite led to a substantial decrease in the temperature of hydrogen evolution, down to room temperature.10 The anthracite was milled with a solvent (cyclohexene) that was dehydrogenated during milling to provide the hydrogen.11 The room temperature hydrogen evolution was accelerated with heating, with the first increase occurring at ∼60 °C and a second increase occurring at ∼100 °C. The material continued to evolve gases when heated in vacuo for up to 4 days.12 High-frequency Raman vibrations at ambient conditions, present only after hydrogenative milling and not in the anthracite precursor, were at the frequency of molecular hydrogen. Observations of nanocrystalline diamond10 and other unusual carbon structures13 further suggested quite unusual transformations. The desorption temperature was significantly less than the hydrogen evolution temperature observed for graphite,1 activated carbon,14 or single-walled nanotubes (SWNTs)9 when milled directly in hydrogen. The hydrogen interactions were atypical of physical adsorption, as there was no associated hydrogen gas phase. Room-temperature hydrogen evolution cannot be attributed to metal hydrides, hydrogen spillover materials, or metals imparted to the carbon during milling.14 The structure of anthracite, recently reviewed by Marzec et al.,15 differs markedly from that of graphite and other engineered carbon materials. In brief, anthracites are the highest rank of coal, classified by the fixed carbon content and the carbon to hydrogen ratio: the carbon content in the organic fraction of an anthracite is in the range of 90-96%, with the bulk of the remaining material made up of hydrogen (∼2%), oxygen * Corresponding author. Phone: 814-863-6256; Fax: 814-865-3248; E-mail:
[email protected]. † Department of Energy & Mineral Engineering. ‡ Materials Research Institute. § EMS Energy Institute. | Department of Chemistry.
( DS) established by the other characterization methods. Summary of Precursor Characterization. In summary, XRD, TPO, and UV Raman all indicate the carbon order is G . DSHT > DS. The ash content of both DS and DSHT is low (1.2 and 1.4 wt %, respectively) and of the same elemental composition. The 1673 K thermal anneal has removed volatile matter, increased the in-plane crystallite dimension, and decreased the cross-link density of the material. Thus, DS represents an anthracite with partially mobile carbon fractions, DSHT represents an anthracite with immobilized carbon fractions, and the most ordered material (G) was used as a control.
Transformations in Carbon Order Induced by Ball Milling. Graphite. Graphite is the most ordered precursor and serves as the baseline for comparison of the effect of precursor carbon order on structural changes during milling. Reactive ball milling in cyclohexene increases the d-spacing of graphite and reduces the crystallite size along the c-axis (Figure 1, Table 1). The 10.9 wt % ash content of mG (Table 1) reflects the expected attrition of the stainless steel milling components during milling. Stainless steel and iron both have XRD reflections around ∼45°, in the range of the [101] peak of graphite. For this reason, the in-plane dimensions (La) are not determined for mG, or any of the milled samples. With the broadening of the [101] peak in mG (due either to particle size reduction or metal addition), the [100] peak is barely distinguishable. An asymmetry in the graphite [002] reflection for mG, previously referred to as γ-phase in the coal literature,21,35 is suggestive of turbostratic carbon. In TPO, the reduction of Tmax with milling (Figure 3) is an effect of the reduced crystallite dimension evident in XRD and the catalytic effect of added metal. mG also has a high temperature shoulder at ∼ 540 °C that is not found in the other milled materials. A fraction (∼8% based on weight loss, Figure 2a) of mG oxidizes below ∼300 °C, and this low-temperature oxidation does not follow ash content of the other materials and is thus attributable to amorphous carbon. The UV Raman of mG is indistinguishable from the graphite precursor, with the same position (1582 cm-1) and breadth (24 cm-1) of the G peak and no indication of a D peak (Figure 3b, Table 1). According to Ferrari and Robertson,29 the D-peak arises from
Effect of Carbon Order During Reactive Ball Milling
J. Phys. Chem. C, Vol. 112, No. 44, 2008 17431
Figure 4. (a) Nitrogen adsorption at 77 K, and corresponding (b) microporosity and (c) mesoporosity of the materials based on DFT. Pore size distribution of precursors is not shown, as the N2 uptake of these materials was minimal.
the breathing motion of sp2 rings and is most evident for infrared and visible excitation, decreasing significantly (to zero) for UV excitation. Nanocrystalline graphite does not give rise to a D peak in UV excitation, thus the absence of a D-peak in UV Raman suggests the material is primarily nanocrystalline graphite and has been reduced in crystallite size without amorphization (Stage 1 in Ferrari and Robertson’s amorphization trajectory). The ∼8 wt% amorphous content (from TPO) is likely masked in UV Raman by the much larger scattering crosssection of nanocyrstalline graphite, which makes up 92% of the sample. Milling of G increases the swelling ratio (Table 2); increased swelling is attributable to amorphous (rather than turbostratic or defected) carbon that has some degree of cross-linking. Helium density of mG decreased slightly, despite an increase in ash content, suggesting certain regions within the amorphous fraction became closed off and inaccessible to helium. The BET surface area of mG is increased ∼10-fold with ball milling
relative to G (Table 2), and the increased surface area is primarily associated with mesopores (Figure 4.) Only ∼0.7% of the total pore volume of mG is attributable to micropores (Table 2). The combined results indicate crystallite size reduction and amorphization. The γ-phase in XRD is suggestive of the introduction of defects and disruption of long-range order. There is no sign of exfoliation in XRD. The low-temperature oxidation in TPO is suggestive that complete amorphization occurs for ∼8 wt % of the material. The combined results for graphite thus suggest that the dominant transformation for graphite is shearing and crystallite size reduction, with a secondary process of amorphization. The resulting material is primarily mesoporous in nature. Demineralized Summit. DS is the least ordered precursor, at the other extreme relative to graphite. Reactive ball milling increases the d-spacing while decreasing the Lc dimension (Table 1). Once again, an increased asymmetry in the [002] peak may
17432 J. Phys. Chem. C, Vol. 112, No. 44, 2008 be a sign of γ-phase carbon21,35(Figure 1). Iron (III) oxide hydroxide is evident in mDS (indicated by *, Figure 1), a likely sign of oxidation during storage. A shift in Tmax with milling (Figure 2, Table 1) is likely due to the catalytic effect of metals added by ball milling. The breadth of the TPO profile for mDS has decreased relative to that of DS, and is suggestive of homogenization of carbon with milling. Ash content indicates the least amount of attrition for all the samples. In UV Raman, the decrease in the G-peak frequency indicates the material has become less “chain-like” and more graphitic.36 The shift in G-peak frequency with milling is the greatest for DS compared to the more ordered precursors. A small decrease in the fwhm for the G-peak is seen for mDS relative to DS: the fwhm (G) is a probe of the structural disordersthe smaller the fwhm (G), the less defects and strain in the sp2-clusters.36 No change is observed in the D/G ratio upon milling. Ball milling increases the swelling ratio of DS from 124% to 150% (Table 2), an indication of increased cross-link density with milling. Yet no net effect is seen for changes in helium density with milling, suggesting equal accessibility to He. Combined, this suggests an alteration of the nature of the chemical bonding that leads to the 3-D networking. SSA remarkably increases 80-fold with ball milling (Figure 4a, Table 2). The 0.073 cc/gc microporosity makes up 23% of the pore volume of mDS, compared to 10% for mDSHT and 0.7% for mG. A large fraction of the µPV of mDS are smaller than 8 Å (Figure 4b). Heat-Treated Demineralized Summit (DSHT). DSHT represents the intermediate case, and is DS partially graphitized by a 1673 K thermal anneal, as discussed above. After milling, mDSHT has a large ash content, 53.1 wt % (Table 1, Figure 2), 5 times higher than any of the other milled samples. A prominent reflection in XRD at 2θ ∼ 45° is likely due to stainless steel, which has a reflection in this range. The reason for high attrition of the milling components for this sample is not clear, but may be due to the high fraction of exposed [102] faces in the DSHT precursor. The relative intensity of the carbon [002] reflection is reduced in XRD as a result of this high metal content, but high resolution of this region (see Supporting Information) indicates the d-spacing of mDSHT has been increased relative to DSHT precursor while Lc remains unchanged (Table 1). The TPO behavior of the DSHT series more closely resembles that of G than that of DS (Figure 2): there is a large shift in Tmax with milling and low-temperature oxidation attributable to amorphous carbon. In UV Raman, the increase in G-peak frequency from DSHT to mDSHT is small (Table 1). A decrease in the fwhm (G) is observed, indicating the ball milling process has reduced defects and strain in mDSHT relative to DSHT.36 As in the other cases, no change is observed in the D/G ratio upon milling. The high ash content of mDSHT is reflected in the swelling ratio of 100% (ash is not expected to swell) and high helium density. Ash content will also affect SSA and pore volume, but an estimate of the SSA and pore volume of the carbon fraction can be obtained by normalizing the per carbon content rather than the total mass. This assumes that the ash content has a small contribution to SSA and porosity, which is a reasonable assumption. Normalizing per carbon content, the SSA of mDSHT is over 38 times that of the DSHT precursor. mDSHT has decreased microporosity relative to that of mDS (Figure 4b) and increased mesoporosity relative to that of mDS and mG (Figure 4c). Comparison of the Order/Structure of Milled Products. As established above, the order of the precursors was G . DSHT
Sakti et al. > DS. It should also be emphasized that DS was the precursor to DSHT. The 1673 K anneal used to convert DS to DSHT immobilized certain carbon fractions and removed volatile matter. With these changes induced by the 1673 K thermal anneal, the behavior of DSHT in the mill most closely resembles that of graphite. Both mDSHT and mG have low-temperature TPO peaks that are attributable to amorphous carbon. Thus, portions of both DSHT and G are fully amorphitized with milling, whereas DS is not. The low-temperature TPO peaks also suggest that more ordered precursors lead to more heterogeneous milled products. Changes in G-peak positioning in UV Raman indicate a change of carbon structure for DS with milling, an effect not observed for G or DSHT. There is some indication by changes in the fwhm (G) that reactive ball milling relieves strain in the carbon materials with inherent disorder. Comparing milled products, sp2 clustering (i.e., increased G-peak frequency) increases as follows: mDS < mDSHT < mG. The trend has not changed compared to that of the carbon precursors prior to ball milling. Comparison of the Order of a 3-D Carbon Network and the Structure of Milled Products. It is interesting to note that the increase in solvent swelling for both mG and mDS suggests an apparent increase in cross-linking density upon ball milling. For mG, this is likely due to the amorphous regions that are evident in TPO. For mDS, no low-temperature oxidation is apparent in TPO, but the structure is clearly turbostratic and disordered. The increased cross-linking density for mDS and mG relative to their precursors may be associated with γ-phase carbon, a region with increased d-spacing between turbostratic layers would be more likely to swell. Both N2 adsorption and helium density primarily probe regions that are disordered or amorphous, as these regions have higher surface area for gas adsorption/interaction. An increase in the N2 BET surface area and porosity is seen upon ball milling for all three samples (Table 2). The total pore volume of all milled materials is comparable (0.2-0.3 cm3/g). Interestingly, the micropore volume (pore diameter < 20 Å, Figure 4b) follows decreasing carbon order of the precursor. The microporosity of mDS is 3.5 times greater than that of mDSHT, and the microporisity of mDSHT is 10 times greater than that of mG. Although TPO indicated that G and DSHT were more susceptible to change with milling, it appears these changes tended to favor creation of mesoporosity and amorphous carbon, whereas the subtle changes seen in TPO for DSHT led to regions of microporosity. The differential pore size distribution (Figure 4b) indicates that the pores of mDS may be in the ultramicropore range, whereas the pores of mDSHT are primarily >10 Å, and the pores of mG are primarily above 20 Å. Disorder in the carbon precursor appears to lead to a more “tightly knit” structure upon reactive ball milling with pores on the molecular scale. There is no correlation between γ-phase carbon in XRD and porosity, as the d-spacing associated with the γ-phase is much smaller than the limit of the smallest pore width detectable by N2 gas adsorption. Previously, ultramicroporosity has been found in graphite milled in a planetary mill in both Ar and H2 atmospheres.37 Hydrogen Evolution. H2 evolutions (m/z ) 2) from DS, mDS, and mG are plotted together in Figure 5. mDSHT was not analyzed because of the high mineral matter content. Also included is the hydrogen evolution of mDS-dry, which is DS milled without added cyclohexene; this sample was added to test the temperature of hydrogen evolution when cyclohexene was not added to the mill. The bar for G represents the
Effect of Carbon Order During Reactive Ball Milling
J. Phys. Chem. C, Vol. 112, No. 44, 2008 17433
Figure 5. Hydrogen evolution for various samples, including controls. The error bars in the figure represent the expected (scaled) drift in the H2 signal based on the G control. Each signal is normalized per initial sample mass.
normalized baseline drift as describe above. It should be stated at the onset that the TDS measurements were conducted over 12 months after milling, thus the hydrogen evolution may be decreased as a result of this prolonged storage time. Hydrogen evolution from the DS precursor originates at ∼600 °C, and ∼400 °C for mDS-dry. This high-temperature (>400 °C) hydrogen evolution for DS and mDS-dry is attributed to gasification of coal with trace moisture, as the TGA-MS is not an entirely sealed system. The onset of the 200 °C decrease in temperature is consistent with a decrease of particle size with milling, which would increase the reactivity of the coal. The H2 evolution from mG is consistent with previous reports of hydrogen evolution from graphite ball milled in hydrogen, which typically have a maxima in the H2 evolution at ∼400 °C.1-9 With complementary neutron and X-ray diffraction,38 Fukunaga et al. attributed this peak at 400 °C for graphite milled in H2 to hydrogen intercalated between graphene layers. A secondary peak at ∼720 °C (990 K) was attributed to hydrogen chemically bound to terminal carbon groups as a result of the hydrogenative ball milling process. Studies by Narayanan and Lueking12 suggest that the sharp peak seen by Orimo et al.1 at 720 °C was due to metals, as Narayanan and Lueking see a peak that is remarkably similar only upon the addition of magnesium. Here, high-temperature H2 evolution is greatest for mDS. On the basis of previous assignments,38 it is likely that the prominent peak at 700 °C for mDS is associated with dehydrogenation of terminal carbons. mDS is expected to have more terminal carbons than the primarily nanocrystalline mG, explaining the increased intensity of high-temperature H2 evolution for mDS relative to that of mG. Low-temperature (i.e., T < ∼300 °C) hydrogen evolution for mG is within the error bar estimated from equipment drift. H2 evolution below 330 °C has not been previously reported for ball milled graphite, nanotubes,9 or other engineered carbon materials.14 The hydrogen evolution from mDS below ∼200 °C is unique for the samples in this study and beyond the error associated with long-term signal drift. Per unit mass, the H2 evolution from mDS is 6-8 times that of mG in the temperature range of 50-200 °C. It does not appear to be associated with hydrogenation of the terminal carbons, as hydrogenated graphite evolves H2 at much higher temperatures.7 The low-temperature hydrogen evolution for mDS does not parallel cyclohexene and
Figure 6. Evolution of cyclohexene (mass 82), benzene (mass 78), and hydrogen (mass 2) shows a correlation for mG (a) but not for mDS (b). Each signal is normalized per initial sample mass. Additional signals are available in the Supporting Information.
benzene evolution as it does for mG (Figure 6): cyclohexene evolved from mDS and mG at temperatures below ∼200 °C. The cyclohexene signal suggests desorption and subsequent evolution of residual cyclohexene from the samples. Benzene evolution for mG paralleled cyclohexene evolution, but lagged behind that of cyclohexene for mDS. This lag for mDS suggests slowed desorption of benzene from the milled coal structure, a likely effect of the increased microporosity. The lag is an indication of dehydrogenation of cyclohexene within the milling process, which was also developed previously.11 The initial increase in hydrogen evolution for mG follows masses 82 and 78, whereas hydrogen evolution for mDS is more complicated. Hydrogen evolution from mDS increases with masses 82 and 78 initially, but continues to increase after masses 82 and 78 subside. This is evidence of hydrogen trapping in mDS due to reactive ball milling in cyclohexene and is discussed further below. Considering all other mass signals tracked (see Supporting Information), there is no clear correlation between the H2 signal and other evolving species. As stated above, the hydrogen evolution was measured approximately 12 months after milling. The lack of correla-
17434 J. Phys. Chem. C, Vol. 112, No. 44, 2008 tion between the H2 signal and that of other gases suggests that the [low-temperature] H2 was formed during the milling process rather than as a byproduct of sample heating. It appears that the H2 evolution from the samples is slowed by the extreme tortuosity of the materials, which is closely related to the ultramicroporosity measured for mDS. Similarly, benzene (the expected decomposition product of cyclohexene dehydrogenation) evolution is slowed for mDS compared to that for mG (see Supporting Information). The low-temperature H2 evolution from mDS is qualitatively similar to that of previous studies of H2 evolution from a Buck Mountain (BMT) anthracite coal milled in cyclohexene,10 although TDS for milled BMT was conducted soon after milling. BMT was generally used without prior demineralization (the exception is when other metals were added to the mill,12,27 in which case BMT was demineralized prior to milling). In addition to unique low-temperature H2 evolution, BMT showed evidence of carbon crystallization to nanocrystalline diamond and unusual high-frequency vibrations in visible Raman.10 The results presented here indicate that the low-temperature H2 evolution is not unique to BMT, although BMT is an anthracite coal, whereas Summit is a semianthracite. One might be tempted to assign the low-temperature hydrogen evolution fully to the ultramicroporous nature of the samples (i.e., mDS); however, milled BMT had a surface area of only 35.5 m2/g12 and no significant microporosity. Clearly the relationship between slowed hydrogen evolution from these samples is more complex than can be probed by simple N2 adsorption at 77 K. It has long been established that N2 diffusion into coals is slowed at 77 K, and N2 surface area may underestimate surface area assessed by CO2 adsorption.30 Furthermore, coals may swell as they imbibe gases that alter their structure and 3-D networking, the most notable example being CO2.30 These observations are expected to be true, if not enhanced, for coals after hydrogenative milling. Furthermore, gases trapped in pores as they are formed are likely inaccessible to gases adsorbed from the gas phase. Summary and Conclusions Precursor order had a substantial effect on the structural changes seen by reactive ball milling in cyclohexene. Milling of graphite (G), the most ordered precursor, led to a heterogeneous structure consisting of both amorphous (TPO) and graphitic carbon (XRD, UV Raman). The graphite crystallites were an order of magnitude smaller (Lc in XRD) after milling. Amorphization led to an increased surface area that was made up of primarily mesoporous carbon that swelled upon the addition of a solvent. The swelling was an indication of regions with a networked, 3-D structure. Hydrogen evolution from milled graphite was similar to that of previous reports, with a maximum in the H2 evolution profile occurring at ∼400 °C. The least ordered precursor, DS, had increased “γ-phase” carbon (XRD), increased clustering or chain-like carbon (UV Raman), increased swelling, and a significantly enhanced N2 adsorption after milling. There was no evidence of lowtemperature oxidation characteristic of truly amorphous carbon in TPO. Ball milling of DS led to the most significant increase in surface area, with significant microporosity for pores smaller than 8 Å. Hydrogen evolution from milled DS was observed at low temperature immediately after heating the sample, despite prolonged storage prior to measurement. The hydrogen evolution from mDS exceeded that of milled graphite by a factor of 6-8 in the temperature range of 50-200 °C.
Sakti et al. Annealing DS prior to milling reduced the effect of milling, likely due to lowering the mobility of carbon fractions within the structure. Similar to graphite, mDSHT had evidence for amorphous carbon in TPO: an increased surface area relative to its precursor that was made up of pores that were both micro- and mesoporous. Annealing the sample prior to ball milling leads to changes that resemble graphite rather than the coal precursor. The results conclusively showed that anthracite behaves differently than graphite in the ball mill. The structural transformations that occur during milling of DS, and previously for BMT anthracite, trap hydrogen for slowed release. The lowtemperature hydrogen evolution is accompanied by a significant ultramicroporsity in the milled DS, but this cannot fully account for the hydrogen evolution, as it was not previously observed for BMT anthracite. The low-temperature hydrogen evolution is likely due to a highly tortuous path within the disordered and tightly knit carbon structure. Acknowledgment. This work was funded through the U.S. Department of Energy University Coal Research Program (DEFG26-06NT42675) and Consortium for Premium Carbon Products from Coal (DEFC2603NT41874, Internal Agreement No. 2875-TPSU-DOE-1874). Supporting Information Available: XRD profile fitting, quantification, Lorentzian peak fits of UV Raman, and additional TDS data is available free of charge via the Internet at http:// pubs.acs.org. References and Notes (1) Orimo, S.; Majer, G.; Fukunaga, T.; Zuttel, A.; Schlapbach, L.; Fujii, H. Appl. Phys. Lett. 1999, 75, 3093–3095. (2) Orimo, S.; Fujii, H.; Matsushima, T.; Ito, K.; Fukunaga, T. Met. Mater. Int. 2000, 6, 601–603. (3) Orimo, S. I.; Matsushima, T.; Fujii, H.; Fukunaga, T.; Majer, G.; Zuttel, A.; Schlapbach, L. Mol. Cryst. Liq. Cryst. 2002, 386, 173–178. (4) Zuttel, A.; Orimo, S. MRS Bull. 2002, 27, 705–711. (5) Fujii, H.; Orimo, S. Physica B 2003, 328, 77–80. (6) Orimo, S.; Zuttel, A.; Schlapbach, L.; Majer, G.; Fukunaga, T.; Fujii, H. J. Alloys Compd. 2003, 356, 716–719. (7) Fukunaga, T.; Itoh, K.; Orimo, S.; Aoki, K. Mater. Sci. Eng., B 2004, 108, 105–114. (8) Hirscher, M.; Becher, M.; Haluska, M.; Quintel, A.; Skakalova, V.; Choi, Y. M.; Dettlaff-Weglikowska, U.; Roth, S.; Stepanek, I.; Bernier, P.; Leonhardt, A.; Fink, J. J. Alloys Compd. 2002, 330, 654–658. (9) Hirscher, M.; Becher, M.; Haluska, M.; von Zeppelin, F.; Chen, X. H.; Dettlaff-Weglikowska, U.; Roth, S. J. Alloys Compd. 2003, 356, 433–437. (10) Lueking, A. D.; Gutierrez, H. R.; Fonseca, D. A.; Narayanan, D. L.; Van Essendelft, D.; Jain, P.; Clifford, C. E. B. J. Am. Chem. Soc. 2006, 128, 7758–7760. (11) Included in the supplementary information of ref 10. (12) Narayanan, D. L.; Lueking, A. D. Carbon 2007, 45, 805–820. (13) Lueking, A. D.; Gutierrez, H. R.; Jain, P.; Van Essendelft, D. T.; Burgess-Clifford, C. E. Carbon 2007, 45, 2297–2306. (14) Ichikawa, T.; Chen, D. M.; Isobe, S.; Gomibuchi, E.; Fujii, H. Mater. Sci. Eng., B 2004, 108, 138–142. (15) Marzec, A. Fuel Process. Technol. 2002, 77, 25–32. (16) Franklin, R. E. Proc. R. Soc. A 1951, 209, 196–218. (17) Iino, M. Fuel Process. Technol. 2000, 62, 89–101. (18) Sanada, Y.; Kumagai, H.; Sasaki, M. Fuel 1994, 73, 840–842. (19) Sanada, Y.; Sasaki, M.; Kumagai, H.; Aizawa, S.; Nishizawa, T.; Mineo, T.; Chiba, T. Fuel 2002, 81, 1397–1402. (20) Bishop, M.; Ward, D. L. Fuel 1958, 37, 191–200. (21) Lu, L.; Sahajwalla, V.; Kong, C.; Harris, D. Carbon 2001, 39, 1821– 1833. (22) Bustin, R. M.; Rouzaud, J. N.; Ross, J. V. Carbon 1995, 33, 679– 691. (23) Liotta, R.; Brons, G.; Isaacs, J. Fuel 1983, 62, 781–791. (24) Green, T. K.; Kovac, J.; Larsen, J. W. Fuel 1984, 63, 935–938. (25) Van Niekerk, D. Ph.D. Thesis, The Pennsylvania State University, University Park, PA, 2008.
Effect of Carbon Order During Reactive Ball Milling (26) Iino, M.; Takanohashi, T.; Ohsuga, H.; Toda, K. Fuel 1988, 67, 1639–1647. (27) Sakti, A. MS Thesis, The Pennsylvania State University, University Park, PA, 2008. (28) Pappano, P. J. M.S. Thesis, The Pennsylvania State University, University Park, PA, 2000. (29) Ferrari, A. C.; Robertson, J. Phys. ReV. B 2001, 64, 075414. (30) Mahajan, O. P. Carbon 1991, 29, 735–742. (31) Polansky, T. S.; Donald, H. J.; Kinney, C. R. Nature 1960, 186, 792–793. (32) Flory, P. J. Principles of Polymer Chemistry; Cornel University Press: Ithaca, NY, 1953.
J. Phys. Chem. C, Vol. 112, No. 44, 2008 17435 (33) Suuberg, E. M.; Otake, Y.; Langner, M. K.; Leung, K. T.; Milosavljevic, I. Energy Fuel 1994, 8, 1247–1262. (34) Takanohashi, T.; Iino, M.; Nishioka, M. Energy Fuel 1995, 9, 788–793. (35) Ergun, S.; Tiensuu, V. H. Fuel 1959, 38, 64–78. (36) Casiraghi, C.; Ferrari, A. C.; Robertson, J. Phys. ReV. B 2005, 72, 085401. (37) Touzik, A.; Hentsche, M.; Wenzel, R.; Hermann, H. J. Alloys Compd. 2006, 426, 272–276. (38) Fukunaga, T.; Itoh, K.; Orimo, S.; Aoki, M.; Fujii, H. J. Alloys Compd. 2001, 327, 224–229.
JP805732X
