Impact of Film Thickness on the Morphology of Mesoporous Carbon

ARTICLE pubs.acs.org/Langmuir

Impact of Film Thickness on the Morphology of Mesoporous Carbon Films Using Organic
Organic Self-Assembly Bryan D. Vogt,*,† Vicki L. Chavez,†,|| Mingzhi Dai,† M. Regina Croda Arreola,† Lingyan Song,†,^ Dan Feng,‡ Dongyuan Zhao,‡ Ginusha M. Perera,§ and Gila E. Stein*,§ †

Chemical Engineering Program, Arizona State University, Tempe, Arizona 85287, United States Department of Chemistry, Fudan University, Shanghai 200433, P. R. China § Department of Chemical and Biomolecular Engineering, University of Houston, Houston, Texas 77204, United States ‡

ABSTRACT: Mesoporous polymer and carbon thin films are prepared by the organic
organic self-assembly of an oligomeric phenolic resin with an amphiphilic triblock copolymer template, Pluronic F127. The ratio of resin to template is selected such that a body-centered cubic (Im3m) mesostructure is formed in the bulk. However, well-ordered mesoporous films are not always obtained for thin films (500 nm) films. The improved stability of orthorhombic mesostructure is not surprising as previously it has been shown for the silica films templated by the same block copolymer, the cubic (orthorhombic) mesostructures have improved thermal stability because of their larger pore walls in comparison to hexagonal (rectangular) structures.34 Recent reports of high-temperature vapor reaction into a preformed template illustrated that monolayers of mesopores can be formed in carbon component, but the authors were unable to obtain the same structure from a solution process.23 Thus, it is unclear that ultrathin films of mesoporous carbon can be produced via an organic
organic self-assembly route, and this question is addressed herein. To assess the morphology present within the films, mesoporous films were characterized by XRD as shown in Figure 2 for films with a 1:0.004 and 1:0.006 phenol to template molar ratios. Thick (>500 nm) mesoporous films with the same phenol/ template ratio have been shown via GISAXS in thick films to exhibit an (010) oriented Fmmm mesostructure.35 From the XRD profiles, only a single diffraction peak is observed that precludes unambiguous assignment of the space group. However, this primary peak position is consistent with the d spacing for the Fmmm mesostructure previously reported. Besides these 5609

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Figure 2. XRD profiles of polymer mesoporous films for phenol:template ratios of (a) 1:0.004 and (b) 1:0.006. For the phenol:template ratio of 0.006, diffraction peaks are observed down to 28 nm thickness, while at a phenol:template ratio of 0.00, there is no sign of a diffraction peak in films as thick as 65 nm. Curves are offset for clarity.

changes, there appears to be some film thickness-dependent behavior that is dependent upon the phenol-to-template ratio. For a ratio of 1:0.004, the primary diffraction peak appears to vanish when the film thickness decreases below 65 nm. This lack of diffractions could be due to a lack of signal as a result of a limited volume of sample being probed, but as shown in Figure 2b, a clear diffraction peak can be obtained from much thinner films with ratios 1:0.006. A diffraction peak is clearly obtained from films as thin as 28 nm for a phenol to F127 ratio of 1:0.006. It is important to note that for film thickness below 100 nm, Kiessig fringes resulting from the interference from the finite film thickness can be observed in some cases. The presence of these fringes indicates that these films are very smooth as the angular divergence from the slits is quite large. The diffraction peak from the ordered structure can be clearly distinguished from the Kiessig fringes by the intensity decay. The interference from the finite film thickness results in a decay as q
4,36 so increased intensity and asymmetric peak are clear indicators of diffraction from the film. This is similar to the observed profiles from neutron reflectivity from block copolymer thin films.37 These results suggest that there is a finite thickness (>65 nm), at least when the ratio of phenol-to-template is 1:0.004, where the ordered mesostructure is not formed or stable after template removal. To further investigate this behavior, the porosity of the films is assessed using ellipsometric porosimetry (EP) as shown in Figure 3. These results are from the mesoporous polymers that have been processed at 350 C as the optical transparency of these films makes the analysis facile; the mesostructured polymer film can be readily transformed into a conductive carbon by reheating at elevated temperatures (>650 C) as can be assessed by the increase in refractive index and orders of magnitude increase in electrical conductivity. The thick films (>100 nm) of both compositions appear to have approximately 25 vol % pores. This porosity is less than commonly reported for mesoporous silica films that are templated by the same surfactant (33%);33 however, there is a significant contraction in the film thickness during template removal.35 For the mesoporous titania, film shrinkage has been shown to be associated with a decrease in total porosity by approximately 0.1.38 Thus, this porosity for the film appears to be reasonable for an ordered mesoporous material. However, as the film thickness is decreased, there is an apparent precipitous drop in the film porosity for both compositions examined. At a ratio of 1:0.004, the porosity appears to begin to decrease for films less than 100 nm. For a 65 nm thick film, the porosity is

Figure 3. Porosity extracted from EP for mesoporous polymer films with phenol:F127 ratios of (9) 1:0.004 and (b) 1:0.006 after pyrolysis at 350 C.

decreased by nearly one-half in comparison to the thick films. This decrease in porosity corresponds with the thickness at which diffraction peaks are no longer observed in XRD. This consistency suggests that an ordered structure cannot be formed using the 1:0.004 ratio for ultrathin films. However, for the films synthesized using a ratio of 1:0.006, the porosity remains constant down to a thickness of approximately 20 nm. The collapse of the mesostructure would result in both a loss of ordering and a decrease in porosity for this film as well, but examining the d spacing for the ordered structure (as determined by XRD) demonstrates that this transition occurs for films that have only 4
5 layers of spheres (n = 4
5). At this thickness, the through-plane diffraction will be weak and thus lack of diffraction does not signify lack of ordering. Moreover, the uncertainty in the optical constants grows as the films become thinner as the wavelength is significantly greater than the optical path length. Thus, it is unclear if these ultrathin films maintain an ordered pore morphology that is observed in the thicker films. To further assess the mesostructure in these ultrathin films with a ratio of 1:0.006, analogous thin films are carbonized at 800 C and examined using TEM and GISAXS. Figure 4 illustrates a TEM cross section of an ordered mesoporous carbon film. The micrograph shows the (101) plane of a face-centered orthorhombic (FCO, Fmmm) mesostructure for n = 7 film after carbonization. This structure is consistent with prior reports for ordered mesoporous carbon films.15,35 However, this film is still within the thickness range where diffraction peaks are observed in XRD. To investigate thinner films, GISAXS is utilized on both the as-made and the carbonized films. Figure 5 illustrates the 5610

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Figure 4. TEM micrograph of a cross section of mesoporous carbon film synthesized using a phenol:template ratio of 1:0.006.

Figure 5. GISAXS profiles for thin films using a phenol:template ratio of 1:0.006 of (a) as-made 17 nm thick film where data are consistent with a HEX monolayer of spheres and (b) as-made 35 nm thick film where data are consistent with HCP bilayer of spheres.

scattering profile for two different ultrathin films synthesized with a ratio of 1:0.006. A film that is initially 17 nm thick after thermopolymerization (Figure 5a) shows two resolved scattering peaks along the in-plane (2Θ) axis. Interestingly, this thinnest film (17 nm as-made) has no apparent structure near the surface as determined from measurements below the critical angle. When the incident angle is larger than the critical angle of the polymeric film, two scattering peaks are resolved at qpar = 0.45 and 0.78 nm
1, respectively. This scattering profile is consistent with in-plane hexagonal symmetry (a = n = 16 nm, peak ratio of 1:(3)1/2). Carbonization of this film at 800 C results in contraction of the film thickness to approximately 5 nm. The scattering from the carbonized film interestingly exhibits a single defined diffraction peak at q = 0.43 nm
1, despite the large increase in contrast from the densification of the carbon and loss of the surfactant to yield pores. The loss of higher-order diffraction peaks suggests disruption of in-plane order during carbonization and prohibits identification of the in-plane symmetry. This is counter to prior studies on thick mesoporous carbon films that suggested uniaxial contraction of the film through the thickness and no distortion inplane.35 When the as-made film thickness is increased to 35 nm (Figure 5b), a single-diffraction peak indicative of lateral ordering is detected from measurements of the film surface. The bulk film is consistent with in-plane hexagonal symmetry (a = n = 16.9 nm, peak ratio of 1:(3)1/2), and the absence of out-of-plane Bragg peaks suggests that there are at most two layers of spheres perpendicular to the substrate. For a film that is 71 nm thick after thermopolymerization (Figure 6), scattering is observed along

Figure 6. As-made 71 nm thick film using a phenol:template ratio of 1:0.006. Data from below (a) and above (b) the critical angle are consistent with coexistence between HCP and FCO phases. The firstorder diffraction peak for HCP is located at 0.42 nm
1; adjacent peaks are associated with FCO symmetry. Locations of Bragg peaks for FCO and HCP are noted in red and black, respectively. (c) For carbonized film, in-plane symmetry is unchanged but the out-of-plane interlayer distance shrinks by over 60%.

both the in-plane (2Θ) and the out-of-plane (R) axes. The out-ofplane Bragg peak position shifts to a larger angle (R) after carbonization, which is consistent with contraction of the film thickness. The appearance of out-of-plane Bragg peaks in both asmade and carbonized samples suggests there are 2 unit cells perpendicular to the substrate, or approximately 5
6 layers of spheres. By careful examination of the GISAXS data, a symmetrybreaking transition is detected that is consistent with the coexistence of HCP and FCO phases through the film thickness. This coexistence is most clearly observed in the line profiles included in Figure 6, where three peaks are detected at qpar = 1.41, 0.42, and 0.45 nm
1. The middle peak is associated with the HCP phase, while the other two are associated with FCO. The HCP phase has lattice parameters of a = n = 17.3 nm, b = (3a)1/2 = 29.9 nm, and c ≈ 23.5 nm. The FCO lattice parameters are a = 18.3 nm, b = 27.9 nm, and c ≈ 22.5 nm; the nearest-neighbor spacing for FCO is n = 16.7 nm, which is nearly identical to the HCP lattice parameter in bilayers. This suggests that the nearest-neighbor distance is approximately conserved through the symmetry-breaking transition. Upon carbonization, the interlayer spacing decreases significantly as c ≈ 8.2 nm (Figure 6c), which is consistent with measurements of the carbonized film thickness by ellipsometry (19 nm). It is difficult to quantify the proportions of HCP to FCO through the film thickness, because peak shapes cannot be reliably determined due to significant overlap. The signal from HCP is about twice as strong as the signal from FCO, which could indicate the HCP phase is dominant or that the HCP phase is much better ordered. If the peak shapes are assumed to be the same between HCP and FCO, this would suggest that the 19 nm thick (as determined from ellipsometry) mesoporous carbon film contains 2/3 HCP and 1/3 FCO. 5611

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Figure 8. Intensity line profile for ultrathick as-made film using a phenol:template ratio of 1:0.006. Ratio of second- to first-order diffraction peaks (along the in-plane axis) is 0.47/0.41 = 1.155, which is consistent with BCC (110) symmetry.

Figure 9. Symmetry transition as a function of the number of layers of spheres using a phenol:template ratio of 1:0.006. As-made and carbonized samples are denoted by red and black symbols, respectively. The limits of in-plane hexagonal (HEX) and BCC (110) symmetry are noted by the dashed lines. Figure 7. As-made and carbonized films that are 420 nm thick using a phenol:template ratio of 1:0.006. The symmetry is FCO, but the inplane symmetry matches the BCC (110) plane with b = (2a)1/2. The out-of-plane interlayer separation is slightly “squashed’” relative to the BCC lattice with dimension c ≈ 0.9(2a)1/2 instead of c = (2a)1/2.

When the film thickness is significantly increased to many layers of spheres, the packing symmetry is FCO for the as-made and carbonized films (Figure 7). This is consistent with prior reports for thick mesoporous carbon films.35 The in-plane lattice parameters are a = 19.0 nm and b = (2a)1/2 = 26.7 nm for both as-made and carbonized films, while the out-of-plane lattice parameter shrinks from c ≈ 23.5 nm to 8.4 nm upon carbonization. The in-plane nearest-neighbor distance is n = 16.4 nm, which is similar to the HEX lattice parameter in ultrathin films. In-plane symmetry in these ultrathick films (Figure 7) is consistent with the BCC (110) plane, the closest-packed plane in the BCC lattice. This consistency is best illustrated by comparisons of simulated and experimental data shown in Figure 7 and the peak positions in noted in Figure 8. The measured and predicted peak positions closely agree with each other. Along the out-of-plane axis, the lattice is “squashed” by about 10% compared with a perfect, (110)oriented BCC lattice (FCO). The analysis of the GISAXS data provides clear evidence for a symmetry-breaking transition that depends on thickness. Figure 9 illustrates how the ratio of the inplane lattice parameters b/a depends upon the number of sphere layers through the film thickness. For the thinnest films, a solely HCP morphology is observed, but this transitions to a mixture of HCP and FCO phases when n = 5
6, and then transforms to a fully FCO morphology in ultrathick (bulk-like) films. The transition from HCP to a mixed HCP/FCO phase in thin films is consistent with previous results for neat block copolymer21 and homopolymer-block copolymer thin films.27

To further assess the distribution of the FCO and HCP phases, the surface morphology of the films was examined using SEM as shown in Figure 10. For the thicker film (with a mixed phase determined from GISAXS), the surface shows a well-defined morphology that is near hexagonally packed, consistent with the (110) orientation of FCO, which is consistent with the GISAXS data that suggests a dominantly (110) oriented FCO surface. For thick mesoporous carbon films, a hexagonally corrugated surface has been reported previously as well.35 Peeling a thick mesoporous film from the substrate shows a square packing at the buried interface, consistent with a (010) oriented FCO phase.35 The surface of the thick mesoporous films adopts a (110) orientation of FCO likely as result of surface wetting that should not strongly depend upon film thickness. It should be noted that there are exceptions, particularly for anisotropic mesostructures where the surface wetting in sub-100 nm films can be altered by the substrate.36 Conversely, the surface of a 10 nm thick film (Figure 10a) shows no evidence for an ordered structure, despite two well-defined diffraction peaks in GISAXS. This ill-defined morphology is observed across the entire surface, and a similar morphology is found for the 5 nm thick film. However, GISAXS from below the critical edge of the thinnest as-made film show no evidence for an ordered surface as well; this featureless SEM micrograph is consistent with the GISAXS data. The lack of a well-defined surface structure in these ultrathin films might be the source for a prior report that monolayer mesoporous carbon films cannot be obtained from organic
organic self-assembly methods;23 one additional complication from this study is the need for the proper phenol:template ratio for the ultrathin films as ordered structures can be formed in thick films for a wider compositional window. It is well established that vastly different mesoporous structures (such as cylindrical and spherical pores) exhibit very different moduli 5612

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Figure 10. SEM images of the surface morphology of mesoporous carbon films using a 1:0.006 ratio that are (a) 10 and (b) 19 nm thick.

at the same void fraction.40 Moreover, collapse of the mesopores in surfactant-templated resol during carbonization has been attributed directly to the mechanical properties of the framework.17,18 The porosity of the noncarbonized films shows a decrease in porosity (Figure 3) for the thinnest films, which is also accompanied by an increase in film contraction upon template removal. As the contraction of the thin carbonized mesoporous films (∼ 71%) discussed here is slightly different from that observed for much thicker films (∼65%),32 the mechanical stability during carbonization may be dependent upon morphology and orientation. As the optical constants for the film are strongly correlated with the film porosity, this difference should be evident from ellipsometry and could explain the apparent decrease in porosity for ultrathin films (Figure 3). However, for ultrathin films, the accuracy of the refractive index can be questionable due to path length considerations. For thick films, the films are opaque and an effective refraction index for the porous carbon material (bulk-like property) is obtained without any film thickness information.35 For intermediate thickness films, the carbon is absorbing but light can still be transmitted and reflected from the substrate back through the film. To improve the accuracy of the fit in this regime, a silicon wafer with a thermal oxide (∼ 1 μm thick) layer is used as the substrate and the ellipsometric angles are measured at multiple incidence angles, which impact the path length of the light. The attenuation of the oscillations in Ψ caused by the thickness of the thermal oxide provides a measure to more accurately assess the absorption coefficient for the mesoporous carbon film. Figure 11 illustrates the incident angular dependence of Ψ for a 220 nm thick mesporous carbon film that utilized a 1:0.006 ratio. The ellipsometric data are fit using a generalized oscillator model that provides a route to ensure Kramers
Kronig consistency in the optical constants for the mesoporous carbon. As shown in Figure 11a, the spectroscopic ellipsometry data cannot be fully fit at all angles and wavelengths assuming a uniform optical constant for the mesoporous carbon. Even allowing for a free fit of n and k, there is limited improvement in the quality of the fit. However, it should be noted that these optical constants are consistent with prior reports for thick mesoporous carbon films, which exhibits electrical conductivity on the order of 20 S/cm after carbonization at 800 C.35,41 However, by introducing a gradient in the optical properties of the film, a significant improvement in the fit quality can be obtained (Figure 11b). This result suggests that the film is not uniform through its thickness, consistent with the GISAXS and SEM data for the 19 nm thick film, but this film is significantly thicker where coexistence is not expected. However, a

Figure 11. Spectroscopic ellipsometric data (O) of mesoporous carbon film that is 220 nm thick using a phenol:template ratio of 1:0.006 that is fit (solid lines) using a generalized Kramers
Kronig-consistent oscillator model with (A) uniform or (B) graded optical properties. A significant improvement in the goodness of fit is found when a graded layer is utilized.

change in orientation of the FCO morphology near the surface of the thick films could impact how the film contracts locally during carbonization and ultimately the porosity near the film surface. The form for the gradient in optical constants must be assumed in the fit; for the best fit shown in Figure 11b, an asymmetric exponential function is utilized. Figure 12 illustrates the refractive index profile for the mesoporous carbon film obtained from this best fit. The near surface of the film exhibits a significantly larger n than the bulk of the film. Further, these data provide some insight into the prior report for extremely low porosity ( 4; this transition is consistent with prior studies involving block copolymers.25,27

’ CONCLUSIONS Ordered mesoporous polymer and carbon thin films have been synthesized by using cooperative self-assembly of resol and Pluronic F127. Well-ordered mesoporous films are formed irrespective of the phenol-to-template ratio in the casting solution if the cast film is thicker than approximately 100 nm. For thinner films, the mesostructure in the porous films can be lost for some compositions. This loss of ordering in ultrathin films is consistent with previous reports from Nishiyama and co-workers on the difficulties in forming ordered mesoporous ultrathin films using a soft templating approach.21 However, the stability of the mesostructure in ultrathin films is found to be correlated strongly with the initial composition of the resol and Pluronic F127. Ordered films as thin as a single monolayer (as determined from GISAXS) can be formed with proper choice of this ratio, which surprisingly leads to the thin walls in the thicker films. This result suggests a complex interplay between interfacial forces in the mesostructural evolution during film formation. A HCP phase is found for n < 4, and a mixed HCP/FCO morphology is observed for slightly thicker films with a transition to only FCO phase in thick films. Careful examination of spectroscopic ellipsometry data also infers a different morphology at the surface of the film in comparison to bulk. ’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected] (B.D.V.); [email protected] (G.E.S.). )

Present Addresses

Present address: Sandia National Laboratory, 1001 University Boulevard SE, Albuquerque, New Mexico 87106, United States.

’ ACKNOWLEDGMENT The authors acknowledge financial support by the National Science Foundation under grant no. CBET-0746664. D.Y.Z. is thankful for financial support from NSF of China. V.L.C. acknowledges support from Sandia National Laboratories. G.M.P. and G.E.S. acknowledge support by the National Science Foundation under grant no. EEC/ECCS-0927147. Use of the APS was supported by the U.S. DOE, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC0206CH11357. We thank Joe Strzalka and Jin Wang for GISAXS support. ’ REFERENCES (1) Yanagisawa, T.; Shimizu, T.; Kuroda, K.; Kato, K. Bull. Chem. Soc. Jpn. 1990, 63, 988. (2) Kresge, C. T.; Leonowicz, M. E.; Roth, W. J.; Vartuli, J. C.; Beck, J. S. Nature 1992, 359, 710–712. (3) Zhao, D. Y.; Feng, J. L.; Huo, Q. S.; Melosh, N.; Fredrickson, G. H.; Chmelka, B. F.; Stucky, G. D. Science 1998, 279, 548–552. (4) Lu, Y. F.; Ganguli, R.; Drewien, C. A.; Anderson, M. T.; Brinker, C. J.; Gong, W. L.; Guo, Y. X.; Soyez, H.; Dunn, B.; Huang, M. H.; Zink, J. I. Nature 1997, 389, 364–368. (5) Brinker, C. J.; Dunphy, D. R. Curr. Opin. Colloid Interface Sci. 2006, 11, 126–132. (6) Ciesla, U.; Schuth, F. Microporous Mesoporous Mater. 1999, 27, 131–149. (7) Yamauchi, Y.; Suzuki, N.; Radhakrishnan, L.; Wang, L. Chem. Rec. 2010, 9, 321–339. (8) Yamauchi, Y.; Kuroda, K. Chem. Asian J. 2008, 3, 664–676. (9) Nishiyama, Y.; Tanaka, S.; Hillhouse, H. W.; Nishiyama, N.; Egashira, Y.; Ueyama, K. Langmuir 2006, 22, 9469–9472. (10) Ryoo, R.; Joo, S. H.; Kruk, M.; Jaroniec, M. Adv. Mater. 2001, 13, 677–681. (11) Lee, J.; Yoon, S.; Hyeon, T.; Oh, S. M.; Kim, K. B. Chem. Commun. 1999, 2177–2178. (12) Hu, Y. S.; Adelhelm, P.; Smarsly, B. M.; Hore, S.; Antonietti, M.; Maier, J. Adv. Funct. Mater. 2007, 17, 1873–1878. (13) Ryoo, R.; Joo, S. H.; Jun, S. J. Phys. Chem. B 1999, 103, 7743–7746. (14) Zhang, F. Q.; Meng, Y.; Gu, D.; Yan, Y.; Chen, Z. X.; Tu, B.; Zhao, D. Y. Chem. Mater. 2006, 18, 5279–5288. (15) Tanaka, S.; Katayama, Y.; Tate, M. P.; Hillhouse, H. W.; Miyake, Y. J. Mater. Chem. 2007, 17, 3639–3645. (16) Tanaka, S.; Nishiyama, N.; Egashira, Y.; Ueyama, K. Chem. Commun. 2005, 2125–2127. (17) Song, L. Y.; Feng, D.; Fredin, N. J.; Yager, K. G.; Jones, R. L.; Wu, Q.; Zhao, D. Y.; Vogt, B. D. ACS Nano 2010, 4, 189–198. (18) Song, L. Y.; Feng, D.; Lee, H. J.; Wang, C. Q.; Wu, Q. Y.; Zhao, D. Y.; Vogt, B. D. J. Phys. Chem. C 2010, 114, 9618–9626. (19) Meng, Y.; Gu, D.; Zhang, F. Q.; Shi, Y. F.; Yang, H. F.; Li, Z.; Yu, C. Z.; Tu, B.; Zhao, D. Y. Angew. Chem., Int. Ed. 2005, 44, 7053–7059. (20) Meng, Y.; Gu, D.; Zhang, F. Q.; Shi, Y. F.; Cheng, L.; Feng, D.; Wu, Z. X.; Chen, Z. X.; Wan, Y.; Stein, A.; Zhao, D. Y. Chem. Mater. 2006, 18, 4447–4464. (21) Zhang, F. Q.; Meng, Y.; Gu, D.; Yan, Y.; Chen, Z. X.; Tu, B.; Zhao, D. Y. Chem. Mater. 2006, 18, 5279–5288. (22) Zhang, F. Q.; Meng, Y.; Gu, D.; Yan, Y.; Yu, C. Z.; Tu, B.; Zhao, D. Y. J. Am. Chem. Soc. 2005, 127, 13508–13509. (23) Jin, J.; Nishiyama, N.; Egashira, Y.; Ueyama, K. Chem. Commun. 2009, 1371–1373. (24) Stein, G. E.; Kramer, E. J.; Li, X. F.; Wang, J. Macromolecules 2007, 40, 2453–2460. 5614

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