PAPER www.rsc.org/materials | Journal of Materials Chemistry

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Silica-free syntheses of hierarchically ordered macroporous polymer andcarbon monoliths with controllable mesoporosity†

Zhiyong Wang, Elizabeth R. Kiesel and Andreas Stein*

Received 18th December 2007, Accepted 27th February 2008

First published as an Advance Article on the web 17th March 2008

DOI: 10.1039/b719489g

Hierarchically ordered macroporous polymer and carbon monoliths with walls containing

face-centered cubic or 2D-hexagonal mesopores were synthesized via a facile dual-templating technique

using poly(methyl methacrylate) (PMMA) colloidal crystals and amphiphilic triblock copolymer

surfactants as templates. A nanocasting step using a silica mold was not required. The as-synthesized

nanoporous structures contain both ordered macropores and mesopores, originating from the colloidal

crystal and surfactants, respectively. The mesostructures could be conveniently controlled by

tuning the concentration of the copolymer surfactant. Starting from the PMMA template, only four

major processing stages (precursor infiltration, solvent removal, thermal curing and carbonization)

were involved to prepare the bimodal porous carbon materials. A two-step thermal curing method was

utilized to improve the robustness of the products. On the basis of nanoindentation measurements, the

carbon products were mechanically more stable than hierarchically porous carbon monoliths

synthesized by nanocasting, and the product with the cubic mesopore structure was even more stable

than 3D-ordered macroporous carbon lacking any templated mesopores in the wall skeleton.

Compared with conventional nanocasting strategies, the current method avoids the use of hazardous

hydrofluoric acid that is required to remove a silica template, and therefore the synthetic procedure is

more environmentally benign.

Introduction

Among the variety of nanostructured materials, carbon with

designed porosity is a promising candidate in applications as

varied as gas separation, water purification, chromatography,

catalysis, hydrogen storage, controlled drug release and

electrodes for electrochemical devices.1 All of these applications

can benefit from better control over the pore structure and

morphology of synthetic carbon materials. In this paper we

report straight-forward and reliable syntheses of nanoporous

carbon and polymer materials with well-controlled morpho-

logies at multiple length scales.

Porous carbons with random mesopores have been known for

a long time, for example in the form of activated carbon.

However, for better access to the interior of the carbonaceous

materials and for applications requiring size-selectivity, a more

ordered mesoporous structure is desirable. Since the pioneering

work of Ryoo et al.,2 interest in ordered mesoporous carbon

has been escalating. Initially, ordered mesoporous carbon was

synthesized by a nanocasting strategy with ordered mesoporous

silica as a hard template and organic compounds, such as

sucrose2 and acetonitrile,3 as carbon precursors. In addition,

a modified method was invented by Hyeon and co-workers in

which block copolymer surfactants were carbonized directly

within mesoporous silica frameworks followed by silica

removal.4 Mesoporous carbon materials with space group

Department of Chemistry, University of Minnesota, Minneapolis, MN55455, USA. E-mail: a-stein@umn.edu

† Electronic supplementary information (ESI) available: Fig. S1, S2 andTable S1. See DOI: 10.1039/b719489g

2194 | J. Mater. Chem., 2008, 18, 2194–2200

symmetries p6mm,5,6 I41/a,2,7 and Ia3d8 have been prepared by

nanocasting methods. The available geometries accessible by

nanocasting are limited to templates (usually silica) with inter-

connected mesoporous channels capable of generating robust,

negative carbon replicas.5 Furthermore, all of the above methods

require either hydrofluoric acid or sodium hydroxide solutions

for silica removal, thus complicating the synthesis.

In order to simplify the production of ordered mesoporous

carbon and to circumvent the use of hazardous chemicals, direct

synthesis techniques were developed. Dai and co-workers first

reported the preparation of hexagonally ordered mesoporous

carbon using resorcinol monomers and a self-assembled copoly-

mer polystyrene-block-poly(4-vinylpyridine) as a mesopore-

directing agent.9 The carbonized product has large mesopores

due to the high molecular weight of the copolymer. Later,

Nishiyama et al. synthesized hexagonally ordered mesoporous

carbon thin films using the inexpensive triblock copolymer,

Pluronic F127, resorcinol-formaldehyde and triethyl ortho-

acetate.10 Recently, a more versatile strategy was reported by

Zhao and co-workers, exploiting the cooperative assembly

between resols and block copolymers to produce ordered

mesoporous carbons with p6mm, Ia3d, Im3m and Fm3m space

groups.11,12 Most of these methods generated carbon samples

as powders or thin films, and there have been very few reports

of direct syntheses of monolithic carbon with ordered meso-

porous structure up to now.13 Moreover, mesoporous carbon

particles prepared by these methods have relatively large external

dimensions (typically over one micrometre), which results in

relatively long diffusion path lengths for guest molecules,

limiting mass-transport.

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One solution to the diffusion limitation is to design a hierarchi-

cally ordered porous structure containing both well-defined

macropores and mesopores. This structure can provide excellent

mass transport through macropores and a large surface area

from mesopores within macropore walls.14 Macropores may be

generated, for example, by colloidal crystal templating to form

a three-dimensionally ordered macroporous (3DOM) struc-

ture.15 Our previous studies have shown that 3DOM carbon,

as an anode for lithium ion batteries, has structural advantages

that would permit high lithiation/delithiation rates.16 Chai

et al. reported that the introduction of secondary mesopores

into a 3DOM carbon skeleton by templating with secondary

silica nanoparticles led to an improved electrochemical activity

when the biporous carbon-product was employed as a catalyst

support in a direct-methanol-fuel-cell.14 Our group synthesized

a hierarchically ordered mesoporous carbon monolith via

nanocasting from a biporous silica template.17 To simplify such

syntheses, two new methods were developed recently. Zhao

and co-workers prepared hierarchically porous carbons from

a dual template approach using silica colloidal crystals and

triblock copolymers as the templates.18 At the same time, our

group reported the synthesis of three-dimensionally ordered

macro-/mesoporous (3DOM/m) carbon monoliths using a

poly(methyl methacrylate) (PMMA) colloidal crystal template

and a triconstituent precursor.19 The as-synthesized carbon

monoliths appeared metallic-yellow opalescent under room light

and had easily controllable dimensions dictated by the PMMA

template. However, both methods required hydrofluoric acid

to remove the silica mold, raising environmental concerns.

Recently a silica-free, direct synthesis route of producing ordered

mesoporous carbon nanocubes was communicated by our

group.20

Here, we report novel direct syntheses of hierarchically porous

phenol-formaldehyde-based (PF-based) polymer and carbon

monoliths with both ordered macropores and mesopores

(3DOM/m PF and 3DOM/m C), controllable mesopore architec-

tures and centimetre-sized external dimensions. The overall

strategy is illustrated in Scheme 1. Only four key stages are

involved in the fabrication of the 3DOM/m products: (1) infiltra-

tion of a solution containing the carbon precursor and a surfac-

tant into a colloidal crystal template, (2) thermal curing, (3)

solvent removal, (4) template removal and/or carbonization.

After the fourth stage, biporous ordered polymer or carbon

Scheme 1 Formation of 3DOM/m PF and 3DOM/m C monoliths. The

syntheses involve four basic steps: (1) infiltration of a precursor solution

containing resol solution, copolymer surfactant and hydrochloric acid

into a PMMA colloidal crystal template, (2) thermal cross-linking of

the resol with surfactant micelles within the void space of the colloidal

crystal, (3) solvent removal under dynamic vacuum, (4) template removal

and carbonization of the composite under an inert atmosphere. A

phenol-formaldehyde-based polymer is produced if the 4th step is carried

out at a lower temperature that eliminates PMMA but does not

carbonize the product.

This journal is ª The Royal Society of Chemistry 2008

monoliths are generated as the final products, depending on

the heating temperature. To the best of our knowledge, this is

the first report of direct syntheses of hierarchically ordered

porous carbon monoliths that completely eliminate any silica

intermediate and the use of hydrofluoric acid. As a result, these

syntheses are safer, easier, cheaper and more environmentally

benign than nanocasting methods employing silica templates.

Experimental

Chemicals

Chemicals used in this experiment were obtained from the

following sources: 2,20-azobis(2-methyl propionamidine)

dihydrochloride initiator (AMPD, 97%), and methyl metha-

crylate monomer (MMA, 99%) were purchased from Aldrich.

Phenol (ACS reagent) and formaldehyde (37% aq. solution)

were from Fisher Scientific. Hydrochloric acid (37%) and sodium

hydroxide were from Mallinckrodt Chemicals. Pluronic� F127

(a difunctional non-ionic ethylene oxide–propylene oxide block

copolymer surfactant, EO97PO69EO97, terminating in primary

hydroxyl groups, average molecular weight 12 600) was received

as a gift from BASF. All chemicals were used without further

purification.

Synthesis of 3DOM/m-c and 3DOM/m-h monoliths

The PMMA colloidal crystal template was prepared by

emulsifier-free emulsion polymerization of MMA at 70 �C with

AMPD as an initiator, as described elsewhere.21 The resulting

PMMA sphere suspension was transferred to a covered glass

crystallization dish and stored for several weeks at room

temperature without agitation. As PMMA spheres sedimented

on the bottom of the container, the water was allowed to

evaporate forming opalescent colloidal crystal pieces. 3DOM/m

PF and 3DOM/m C monoliths with two different mesoporous

structures were prepared in a direct dual-templating method. A

50 wt% phenol-formaldehyde (PF) resol solution in ethanol

(pH �7) was first synthesized according to an established

method.12 In a typical synthesis, 2.0 g PF sol was mixed with

1.0 g of 0.2 M HCl(aq) in a glass vial. Next, 0.5 g (for cubic

3DOM/m-c) or 1.0 g (for hexagonal 3DOM/m-h) of F127 was

added to this clear mixture and stirred for at least 6 h at room

temperature to form a homogeneous, yellow solution. PMMA

monoliths were infiltrated with the above precursor solution

inside a capped container. Care was taken that the solution level

was not above the PMMA pieces. After excess solution was

wiped off the monoliths, these pieces were placed in a vacuum

oven for 3–6 h at 50–60 �C to remove the solvent at reduced

pressure (< 0.5 mmHg). Cross-linking was achieved by thermally

curing the precursor-infiltrated monoliths, first at 100 �C for 24 h

in capped plastic bottles and then at 140 �C for 24 h within

loosely covered Petri dishes. Phenolic resin monoliths were

obtained after heating at 450 �C for 3 h with a heating rate of

1 �C min�1 under flowing N2 (1 L min�1). These samples were

denoted as 3DOM/m-cPF (cubic phase) and 3DOM/m-hPF

(hexagonal phase). Carbon monoliths were generated by carbo-

nizing the thermally-cured phenolic resin–PMMA composite

under flowing N2 (1 L min�1) at 450 �C for 3 h and then at

900 �C for another 2 h with a heating rate of 1 �C min�1. These

J. Mater. Chem., 2008, 18, 2194–2200 | 2195

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samples were denoted as 3DOM/m-cC and 3DOM/m-hC,

respectively.

Characterization

C, H, O elemental analyses (EA) of the samples were performed

by Atlantic Microlab Inc. of Norcross, GA. Scanning electron

microscopy (SEM) images were obtained with a JEOL 6700

scanning electron microscope at an accelerating voltage of

5 kV. Samples were ground to a powder and then mounted on

an aluminium stub with conductive carbon tape. Transmission

electron microscopy (TEM) was conducted on a JEOL model

JEM-1210 electron microscope operated at 120 kV. All samples

were crushed to a powder and sonicated in ethanol for 1 h before

TEM analysis. They were drop-dried from acetone onto TEM

grids. Small-angle X-ray scattering (SAXS) data were acquired

on a Rigaku RU-200BVH 2D SAXS instrument using

a 12 kW-rotating anode with a Cu source and a Siemens

Hi-Star multi-wire area detector. The sample-to-detector

distance was 108 cm, exposure time 10 minutes for each sample

and the step interval 0.008� 2q. All samples were ground into fine

powders before running tests. Nitrogen-sorption measurements

were performed on a Quantachrome Autosorb-1 Automated

Gas Sorption System. Samples were degassed at 0.003 mmHg

for at least 12 h at 200 �C. Specific surface areas were calculated

by the Brunauer–Emmett–Teller (BET) method, and pore sizes

and volumes were estimated from pore size distribution curves

from the desorption branches of the isotherms using the

Barrett–Joyner–Halenda (BJH) method. Pore volumes were

taken at the P/P0 ¼ 0.995 single point. Depth-sensing indenta-

tion (DSI) experiments were performed using an instrumented

indenter (Nano Indenter XP, MTS Systems Corporation) with

a diamond Berkovich indentation tip. The maximum load for

each sample was 15–16 mN. Multiple load-unload cycles were

conducted for each sample to examine the reproducibility. To

evaluate the influence of mesoporosity on the strength of

monoliths, a 3DOM carbon monolith derived from a resorcinol-

formaldehyde precursor (3DOM RFC) without mesoporosity

inside macropore walls was prepared according to our previously

developed method.16

Results and discussion

Precursor conditions

As a starting point, a resol precursor composition used in

previous reports of mesoporous carbon films was employed.12

To explore the mesophases attainable in the confinement of

a PMMA colloidal crystal, the surfactant concentration was

varied. On the basis of the analyses described below, hierarchi-

cally ordered polymer and carbon monoliths with an fcc meso-

structure (3DOM/m-c) were synthesized at a low surfactant

concentration in the precursor solution with molar compositions

of 1.0 phenol, 2.0 formaldehyde, 6.1–8.6 � 10�3 F127, 3.1 � 10�2

HCl, 8.5 H2O, 3.3 EtOH. A mesostructure phase transformation

from cubic to 2D-hexagonal was observed when the concentra-

tion of F127 was increased in the precursor solution to molar

compositions of 1.0 phenol, 2.0 formaldehyde, 1.2–1.8 � 10�2

F127, 3.1 � 10�2 HCl, 8.5 H2O, 3.3 EtOH for 3DOM/m-h.

Here the presence of hydrochloric acid was essential for the

2196 | J. Mater. Chem., 2008, 18, 2194–2200

dissolution of F127, probably due to the enhancement of the

interactions between the protic solvents and the oxyethylene

segments in F127. To produce highly ordered macro- and

mesoporous frameworks, it was also vital to remove solvents

(especially ethanol, which is known to retard the formation of

an ordered mesostructure)22 before the thermal curing step. A

material with completely disordered mesopores and macropores

was obtained without this solvent-removal process.

Thermal curing

A two-step thermal curing procedure was employed. The first

step was conducted at 100 �C for 24 h within a capped plastic

container to solidify and pre-cross-link the resol. This tempera-

ture was selected for the following reasons. At temperatures

lower than 100 �C, cross-linking of resol is too slow; however,

excessively high temperatures destroy the surfactant micelles

and lead to macroscopic phase separation between F127 and

the resol. Moreover, temperatures higher than the glass transi-

tion temperature deform the PMMA colloidal crystal (Tg of

PMMA is ca. 105 �C) because the composite is insufficiently

hardened at the early processing stage. Heating in a closed

container is very important during this first step since a sufficient

amount of HCl is required to achieve a high-degree of cross-

linking, and HCl would be slowly lost in an open container.

After this polymerization step, the monolith was orange. The

extent of cross-linking of the phenolic resin is related to its color:

the higher the cross-linking level, the darker the color. Poorly

cross-linked phenolic resin normally gives a pale-yellow color.

The second step involves curing the composite at 140 �C. This

is intended to further harden the resin and avoid corruption of

the mesostructure during the decomposition of the surfactant

and the subsequent carbonization step. This temperature corres-

ponds to an empirically established optimum. Heating tempera-

tures lower than 140 �C did not influence the cross-linking and

mesopore order significantly (compared with curing at 100 �C);

temperatures higher than 140 �C produced samples with closed

mesopores that were not accessible in nitrogen sorption experi-

ments. After this step, the monoliths were dark-brown and

appeared opalescent. It is worth noting that the periodic

structure of the PMMA colloidal crystal was maintained even

though the temperature was far beyond the Tg of PMMA at

this second step. The PMMA spheres were restrained within

a hard phenolic resin framework due to the pre-solidification

from the first step.

Template removal and carbonization

After the two-step curing, the PMMA–phenolic resin composite

was heated under flowing nitrogen at 450 �C to decompose the

copolymer surfactant and the PMMA template to yield

3DOM/m phenolic resin monoliths. When the composite was

heated up to 900 �C, the polymer was converted to a glassy

carbon phase. As-synthesized 3DOM/m monoliths had typical

dimensions of about 0.6 � 0.5 � 0.3 cm3 (Fig. 1). 3DOM/

m-cPF and 3DOM/m-hPF monoliths were dark brown whereas

3DOM/m-cC and 3DOM/m-hC monoliths were black. Cross-

sections of these monolithic samples appeared opalescent as

a result of the periodic structure with repeat lengths on the order

This journal is ª The Royal Society of Chemistry 2008

Table 1 Elemental analysis results of phenolic resin and carbon mono-liths

SampleElement content/wt %

FormulaC H O Total

3DOM/m-cPF 86.09 3.98 8.05 98.12 C14.3H7.8O3DOM/m-cC 94.79 0.22 2.34 97.35 C54.1H1.5O3DOM/m-hPF 86.98 3.96 6.92 97.86 C16.7H9.1O3DOM/m-hC 94.37 0.21 2.38 96.96 C52.7H1.4O

Fig. 1 Photograph of 3DOM/m monoliths.

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of the wavelength of visible light. Phenolic resin samples

(3DOM/m-c and 3DOM/m-h) contained hydrogen and oxygen

in high concentrations (Table 1). After carbonization at 900�C, both the oxygen and hydrogen contents decreased. The total

percentages of C, H and O analyzed were slightly less than

100%, which could be due to some inorganic residues from the

reactants.

Structural analysis

Scanning electron microscopy images revealed that all carbo-

nized monoliths had an ordered macropore structure (Fig. 2).

On the basis of SEM images, the PMMA spheres used to template

3DOM/m C monoliths had diameters of 416 � 11 nm.19

After removal of the PMMA template by heat treatment at

Fig. 2 SEM images of carbon monoliths: 3DOM/m-cPF (A), 3DOM/

m-cC (B), 3DOM/m-hPF (C) and 3DOM/m-hC (D).

This journal is ª The Royal Society of Chemistry 2008

450 �C, a well-interconnected framework composed of phenolic

resin was formed (Fig. 2A and 2C). The macropore diameters

of 3DOM/m-cPF and 3DOM/m-hPF were 394 � 7 and 404 �11 nm respectively (Table 2), indicating relatively little shrinkage

after removal of the PMMA spheres (3–5%). After carbonization

at 900 �C under nitrogen, macropore diameters further decreased

to 342 � 6 and 348 � 7 nm for 3DOM/m-cC and 3DOM/m-hC,

respectively. At this stage, the shrinkage of the 3DOM framework

with respect to PMMA was significant (16–18%). The large

shrinkage might be due to the loss of low-molecular-weight

components during pyrolysis at high temperatures. Although

the amount of shrinkage of 3DOM/m-cC and 3DOM/m-hC was

greater than for a 3DOM/m C monolith prepared by a triconsti-

tuent method using the same PMMA template (12%),19 macro-

pore window sizes of these samples were larger (ca. 150 nm for

3DOM/m-cC vs. 130 nm for 3DOM/m C synthesized via the

triconstituent method). The big windows could be the result of

high temperature annealing of PMMA during the second curing

step which coarsens the necks between adjacent PMMA spheres.

The wider windows can facilitate mass transport if 3DOM/m C is

used for applications such as catalysis and energy storage.

The interior structure of the 3DOM/m-c samples was analyzed

by TEM (Fig. 3 and ESI†, Fig. S1). Regularly patterned, sphe-

rical mesopores were observed across the entire polymer and

carbon samples. Mesopores were well ordered within octahedral

and tetrahedral centers of the 3DOM skeleton (i.e., centers

corresponding to octahedral and tetrahedral holes in the original

colloidal crystal template). The mesostructure became slightly

less ordered near the macropore wall surface. These observations

suggest that interactions between surfactant molecules and the

PMMA surface, as well as confinement effects, influenced the

alignment of F127 micelles. Based on the pore geometries and

pore spacings apparent in the TEM images along the [100],

[110] and [111] viewing directions, a cubic structure correspond-

ing to the Fm3m space group is proposed for mesopores in the

3DOM/m-c samples (see also ESI†, Fig. S2). The orientation

of the mesostructure dovetails with that of the 3DOM structure,

as we also observed in mesoporous carbon nanoparticles.20

TEM images of 3DOM/m-h samples revealed significantly

different mesostructures. 2D-hexagonally ordered mesopores

were observed along the [100] axis of the 3DOM framework

(Fig. 3 and ESI†, Fig. S1). Within small domains, the orientation

of cylindrical mesopores correlated with that in adjacent octa-

hedral centers, but such correlations did not extend throughout

the complete sample. In addition, some distortion was observed

around the core region of mesopore bundles. When viewing the

sample through the [110] direction of the 3DOM structure, it was

noticeable that mesoporous channels followed the curvature of

macropore wall surface. This is additional indirect evidence

that the self-assembly of F127 micelles was influenced by the

PMMA sphere surfaces. A similar phenomenon was reported

in our previous work on 3DOM/m silica monoliths using

Pluronic P123 as mesopore-generating agent.23

Small-angle X-ray scattering was used to confirm the mesopore

symmetries of the 3DOM/m monoliths (Fig. 4). For 3DOM/

m-cPF, a strong scattering peak at 0.80� 2q can be assigned to

the reflection from (200) planes in the cubic mesopore structure

(Fig. 4A). The peak is relatively broad, consistent with small

ordering domains in the mesostructure and in agreement with

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Table 2 Dimensional data from SEM, SAXS and nitrogen sorption

Sample name Macropore size/nma Mesopore size/nmb Mesopore wall thickness/nmc SBET/m2 g�1 Vpore/cm3 g�1 Vmicropore (%)

3DOM/m-cPF 394 � 7 3.0 12.6 376 0.29 273DOM/m-cC 342 � 6 2.7 10.9 505 0.37 163DOM/m-hPF 404 � 11 3.2 8.7 380 0.31 273DOM/m-hC 348 � 7 3.0 7.4 464 0.34 14

a Measured from SEM images. b Measured from the desorption branches of the pore distribution curves using the BJH method. c The mesopore wallthickness (t) is calculated from the pore spacing (dhkl; from SAXS data) and mesopore diameters (D; from nitrogen sorption):

t ¼ O2d200 � D for 3DOM/m-c (1)

t ¼ (2/O3)d100 – D for 3DOM/m-h (2)

Fig. 3 TEM images of 3DOM/m-cC (A, B and C) and 3DOM/m-hC (D, E and F) viewed along (A) [100], (B) [110] and (C) [111] axes of the macro-

porous structure. Corresponding images for 3DOM/m-cPF and 3DOM/m-hPF viewed from the above three directions are shown in the ESI†, Fig. S1.

Fig. 4 SAXS patterns of (A) 3DOM/m-c and (B) 3DOM/m-h monoliths

heated at 450 �C (PF samples) and 900 �C (carbon samples). The inten-

sity of the data trace for 3DOM/m-cC within the range of higher order

peaks was amplified by a factor of five for clarity. Vertical lines above

the trace for 3DOM/m-cC correspond to calculated positions for the

(111), (200), (220) and (222) reflections for 3DOM/m-cC, using a cubic

cell with parameter a ¼ 19.1 nm.

2198 | J. Mater. Chem., 2008, 18, 2194–2200

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the TEM observations that show ordered domains of ca. 150 nm

(ESI†, Fig. S1). On the basis of the d-spacings of the (200) peak

and weaker (220) and (222) peaks, the parameter a of the cubic

unit cell is 22.3� 0.4 nm, close to that estimated from TEM images

(21.4 � 1.6 nm). After carbonization at 900 �C, all diffraction

peaks shift to higher angles due to shrinkage of the mesostructure,

resulting in a cubic cell with parameter a ¼ 19.1 � 0.2 nm

(compared to a ¼ 18.9 � 1.0 nm estimated from TEM images).

A (111) peak is observed only as a weak shoulder, perhaps due

to the strong curvature of the mesopore domains along the [111]

direction. 3DOM/m-hPF shows a major SAXS peak at 0.86� 2q

(Fig. 4B), which is attributed to scattering from the (100) planes

in a 2D-hexagonal mesopore structure with a ¼ 11.9 nm (TEM

estimate: a ¼ 12.7 � 0.3 nm). After carbonization at 900 �C, the

(100) peak shifts to 0.98� 2q for 3DOM/m-hC. The a value

estimated from the (100) peak is 10.4 nm and the corresponding

separation of d100 planes is 9.0 nm, consistent with the TEM

observations (a ¼ 10.8 � 0.4 nm, d100 ¼ 9.4 � 0.3 nm). Because

of the small domains and channel distortions observed in the

This journal is ª The Royal Society of Chemistry 2008

Fig. 5 Results from nitrogen sorption experiments. (A) and (C) are

isotherms while (B) and (D) are pore size distribution curves calculated

by BJH method from the desorption branches. The symbols (hollow circle,

solid circle, hollow triangle and solid triangle) represent data for 3DOM/

m-cPF, 3DOM/m-cC, 3DOM/m-hPF and 3DOM/m-hC, respectively.

Fig. 6 (A) Indentation load–displacement (P–h) responses and (B) plots

of the indentation contact response in pressure-volume (p–V) space for

indentations made for a maximum load of 15–16 mN. Data traces in

red, orange, blue and green colors correspond to 3DOM/m-cC,

3DOM/m-hC, 3DOM RFC and 3DOM/m-nC, respectively. Here

3DOM/m-nC denotes the 3DOM/m carbon monolith prepared by our

previously reported nanocasting method. Data for 3DOM RFC and

3DOM/m-nC were both adopted from ref. 17.

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TEM images, the SAXS patterns for the 3DOM/m-hPF and

3DOM/m-hC samples only approximate those for a well-ordered

2D hexagonal mesostructure. Therefore, the low intensity peaks

between 1.3 and 2.0� 2q were not indexed.

Nitrogen sorption experiments were performed to evaluate the

overall porosity of the polymer and carbon samples (Fig. 5). All

these samples showed type-IV isotherms with narrow hysteresis

loops in the P/P0 range of 0.3–0.8, consistent with the meso-

porous nature of the skeletal walls (Fig. 5A and 5C).24 However,

isotherm shapes are different for the polymer and carbon

samples. The hysteresis loops of 3DOM/m-cPF and 3DOM/

m-hPF do not close, even at very low P/P0 values. This behavior

may be related to the ability of the polymers to swell in liquid

nitrogen during the adsorption measurements.25 This explana-

tion is supported by the observation that the polymer samples

contained almost twice as much microporosity as the carbon

counterparts (Table 2). 3DOM/m-cPF and 3DOM/m-hPF had

BET surface areas of 376 and 380 m2 g�1, respectively, and these

values increased to 505 and 464 m2 g�1 after carbonization at

900 �C under flowing nitrogen. All the 3DOM/m-c and

3DOM/m-h samples exhibited very narrow pore size distribu-

tions with mesopore diameters of ca. 3 nm, and this value

decreased slightly after the transformation from polymer to

carbon (Table 2). The pore size distribution curves show a steep

rise at pore diameters below 2 nm, which indicates that signifi-

cant fractions of micropores are present (Fig. 5B and 5D). The

total pore volumes for 3DOM/m-cC and 3DOM/m-hC are

comparable (0.37 and 0.34 cm3 g�1), and these values are smaller

than for 3DOM/m carbon monoliths we prepared previously via

nanocasting or from triconstituent precursors.17,19

Mechanical properties

The 3DOM/m-cC and 3DOM/m-hC samples were found to be

much stronger than other 3DOM/m carbon monoliths with

hierarchical porosity we reported before,17,19 and they can all

This journal is ª The Royal Society of Chemistry 2008

be handled easily without breakage. To further evaluate their

mechanical strength, depth-sensing indentation (DSI) was

employed. A load was applied on monoliths between 0 and

16 mN while displacements of the indenter tip were recorded

in multi-cycle tests (Fig. 6A). At the same load, the displace-

ments of 3DOM/m-cC with cubic mesopores were smaller than

those reported for 3DOM RF carbon without designed

mesoporosity (3DOM RFC).17 In spite of the hierarchical pore

structure, the dual-templated product displayed a higher

resistance toward deformation than 3DOM RFC prepared only

with a colloidal crystal template. This behavior contrasts with

that of 3DOM/m-hC with 2D hexagonal mesopores. Although

for this sample the reproducibility of the load displacement curves

in multiple test cycles was poor, all curves indicated that 3DOM/

m-hC was less rigid than the other two samples. All three samples

were significantly stronger than 3DOM/m-nC prepared via

a nanocasting route.22 It is difficult to obtain modulus and

hardness values directly from the standard load-displacement

traces in Fig. 6A due to the non-elastic and non-volume-

conserving responses of porous materials.26,27 However, a

material parameter could be derived from indentation pressure

vs. indentation volume plots to compare the ‘‘strengths’’ of these

samples. Fig. 6B shows pressure–volume (p–V) traces calculated

from the typical P–h traces according to a published method.26

Based on these plots, the strengths of the monoliths followed

the sequence 3DOM/m-cC > 3DOM RFC > 3DOM/m-hC >

3DOM/m-nC. It is interesting that the more porous 3DOM/

m-cC monolith was stronger than the 3DOM RFC monolith

without secondary mesopores in the walls. This could be a

function of multiple parameters, including differences in the

secondary curing temperature, precursor type and catalyst used

during the syntheses of these two samples. The 3DOM/m-hC

monolith was not as strong as 3DOM/m-cC since the former

was built up with much thinner mesopore walls (Table 2). In

addition, the macroscopic strength of 3DOM/m monoliths could

also relate to their internal nanostructures. 3DOM/m-cC was

composed of a three-dimensional, interconnected mesoporous

framework while 3DOM/m-hC had a two-dimensional, cylin-

drical mesopore structure. Therefore, the internal structure of

3DOM/m-hC was more deformable and consequently it was

not as strong as that of 3DOM RFC.

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Conclusions

Hierarchical polymer and carbon monoliths with ordered macro-

and mesopores were synthesized via a facile dual-templating

technique involving colloidal crystal and block copolymer

templates. The mesostructure could be easily controlled by tuning

the concentration of the block copolymer. A two-step thermal

curing strategy was adopted to achieve a highly cross-linked

mesostructure composed of phenolic resin and block copolymer.

This technique ensured the formation of robust mesopore walls

containing phenolic resin that survived during the decomposition

of the block copolymer template. The ordered mesoporous

phenolic resin was subsequently transformed to mesoporous

carbon after heat treatment at high temperatures under an inert

atmosphere. Unlike the commonly encountered evaporation-

induced self-assembly pathway during the fabrication of ordered

mesoporous carbon films, formation of ordered mesostructures

within the voids of colloidal crystal template followed quite

a different pathway. The growth of mesopores was significantly

influenced by the confinement effect of the colloidal crystal

template. Both spherical (3DOM/m-c) and cylindrical (3DOM/

m-h) mesopores were aligned parallel to the surface of PMMA

spheres, and therefore the obtained mesostructures exhibited

apparent curvatures near the surface of macropore walls. This

synthesis has several advantages compared with reported strate-

gies so far: (1) it circumvents the use of hazardous hydrofluoric

acid, which is employed in nanocasting methods, making the

current method more environmentally benign; (2) it generates

polymer and carbon samples with tunable mesostructures (cubic

and 2D-hexagonal); (3) it produces carbon monoliths with

relatively high mechanical strength comparable to macroporous

carbon without pre-defined mesoporosity. Porous carbon

monoliths with the above mentioned properties make them ideal

candidates as guest–host substrates, which could find applica-

tions in catalysis, adsorption, separation, sensing, energy storage

and conversion.

Acknowledgements

This research was supported in part by the Office of Naval

Research (grant number N00014-07-1-0608), the Petroleum

Research Foundation administered by the American Chemical

Society (ACS-PRF Grant No. 42751-AC10), and the National

Science Foundation (DMR-0704312). Portions of this work

were carried out at the University of Minnesota I.T. Characteri-

zation Facility, which receives partial support from the MRSEC

program of the NSF (DMR-0212302). Z. Wang thanks the

University of Minnesota Graduate School for a Doctoral Disser-

tation Fellowship. The authors thank Prof. M. Tsapatsis for the

use of the nitrogen sorptometer and Dr J. Nelson for DSI

measurements.

2200 | J. Mater. Chem., 2008, 18, 2194–2200

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