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Carbon Aerogels


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Carbon aerogels are a unique class of high-surface-area materials derived by sol–gel chemistry. Their high mass-specific surface area and electrical conductivity, environmental compatibility, and chemical inertness make them very promising materials for many applications, such as energy storage, catalysis, sorbents, and desalination. Since the first carbon aerogels were made via pyrolysis of resorcinol-formaldehyde-based organic aerogels, in the late 1980s, the field has really grown. Recently, in addition to RF-derived amorphous carbon aerogels, several other carbon allotropes have been realized in aerogel form: carbon nanotubes, graphene, and diamond. Furthermore, use of the carbon-based aerogels as a platform for making polymer composites has produced order of magnitude improvements in the polymer’s conductive and mechanical properties. Finally, functionalization of these new carbon aerogels via surface engineering has led to a host of interesting composite aerogels that could make aerogels promising candidates for an even wider array of applications. In this chapter, we will present recent work covering the novel synthesis of CNT, graphene, and composite aerogels, as well as their performance in a number of applications.
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Carbon Aerogels
Marcus A. Worsley and Theodore F. Baumann
Introduction ....................................................................................... 2
Carbon Nanotube-Based Aerogels and Xerogels . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . 4
Carbon Nanotube-Based Aerogels ............................................................ 4
Carbon Nanotube-Based Xerogels ............................................................ 7
Graphene-Based Aerogels and Xerogels ......................................................... 17
Graphene Aerogels via Resorcinol-Formaldehyde Solgel Chemistry ...................... 18
Graphene Aerogels from Pure Graphene Oxide Suspensions . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . 24
Graphene Xerogels ............................................................................ 28
Conclusion .. . . . ................................................................................... 32
References .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . . . . . . 32
Carbon aerogels are a unique class of high-surface-area materials derived by
solgel chemistry. Their high mass-specic surface area and electrical conduc-
tivity, environmental compatibility, and chemical inertness make them very
promising materials for many applications, such as energy storage, catalysis,
sorbents, and desalination. Since the rst carbon aerogels were made via pyrol-
ysis of resorcinol-formaldehyde-based organic aerogels, in the late 1980s, the
eld has really grown. Recently, in addition to RF-derived amorphous carbon
aerogels, several other carbon allotropes have been realized in aerogel form:
carbon nanotubes, graphene, and diamond. Furthermore, use of the carbon-
based aerogels as a platform for making polymer composites has produced
order of magnitude improvements in the polymers conductive and mechanical
properties. Finally, functionalization of these new carbon aerogels via surface
M.A. Worsley (*)T.F. Baumann (*)
Physical and Life Sciences Directorate, Lawrence Livermore National Laboratory, Livermore, CA,
#Springer International Publishing Switzerland 2016
L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology,
DOI 10.1007/978-3-319-19454-7_90-1
engineering has led to a host of interesting composite aerogels that could make
aerogels promising candidates for an even wider array of applications. In this
chapter, we will present recent work covering the novel synthesis of CNT,
graphene, and composite aerogels, as well as their performance in a number of
Carbon aerogels, invented by Dr. Pekala (Pekala et al. 1990), possess several unique
properties that make them desirable for a number of technologies including energy
storage, catalysis, ltration, and actuators. First, carbon is a fairly light element, so
materials made from it have the potential to be very low density. For example,
though silica aerogels held the title of worlds lightest materialfor a long time
at ~1 mg/cm
, recently, carbon-based aerogels have shattered that record with a
density of less than 200 μg/cm
(Sun et al. 2013). Carbon-based aerogels can also be
engineered to have ultrahigh surface areas. In fact, carbon nanotube (Worsley
et al. 2013) (CNT) and graphene (Worsley et al. 2012) aerogels routinely exhibit
BET surface areas in excess of 500 and 1000 m
/g, respectively, while activated
carbon aerogels (Baumann et al. 2008) have reported values in excess of 3000 m
The sp
bonding prevalent in carbon-based aerogels affords good electrical conduc-
tivity and mechanical properties (e.g., modulus, compressibility) superior to those of
inorganic aerogels at equivalent densities. Finally, recent advances in aerogel syn-
thesis have afforded carbon aerogels that cover nearly all of the known carbon
allotropes, including not only the CNT and graphene aerogels mentioned earlier
but also diamond aerogels. Diamond aerogels, though not discussed in this chapter,
offer access to yet another set of unique optical, electronic, and thermal properties
(Pauzauskie et al. 2011). With such extraordinary and diverse properties, carbon-
based aerogels continue to be the subject of intense research more than a quarter of a
century after their introduction.
The methods of carbon-based aerogel synthesis have expanded signicantly since
the rst carbon aerogels were created through the carbonization of resorcinol-
formaldehyde-derived organic aerogels. Though carbonization of organic aerogels
continues to be a prevalent technique, many methods now use suspensions of highly
-hybridized carbon nanomaterials (e.g., CNTs, graphene, or graphene oxide
(GO)) in the synthesis process. These methods include the use of freeze-drying of
carbon suspensions (Wang and Ellsworth 2009; Bryning et al. 2007) to create
carbon-based aerogels held together primarily by van der Waals forces as well as
solgel processes (Worsley et al. 2009a,2010) that give covalently cross-linked
aerogels (Fig. 1). In the latter case, the concentration of sp
-hybridized carbon in
suspension and the organic gel precursor (e.g., % RF solids) can be tuned such that
the organic particles primarily nucleate on the surface of the sp
-hybridized carbons.
As such, these organic particles mainly serve to covalently cross-link the sp
hybridized carbons with highly conductive, mechanically robust sp
carbon bridges
2 M.A. Worsley and T.F. Baumann
after carbonization. These aerogels no longer consist of a three-dimensional network
of amorphous carbon nanoparticles, but three-dimensional networks of CNTs or
graphene. CNT and graphene aerogels have shown signicant enhancements in
transport and mechanical properties compared to traditional carbon aerogels due to
the crystalline nature of the primary carbon units. In fact, due to the extraordinary
properties of graphene aerogels in particular, research on these materials has
increased dramatically since the rst report in 2009 (Fig. 2).
The main objective of this chapter is to provide an overview on the preparation,
properties, and applications of CNT and graphene aerogels and xerogels. Recent
interests in high-surface-area materials for energy and environmental applications
reect the volume of the descriptions. Though several methods have been attempted
to prepare CNT and graphene aerogels, various covalently cross-linked versions
prepared via solgel methods are chosen for the examples. For information on
traditional carbon aerogels, readers are directed to the excellent review by
Al-Muhtaseb and Ritter (2003).
Fig. 1 Schematic showing the solgel synthesis of graphene-based aerogels
Fig. 2 Graph showing the number of graphene aerogel journal articles published per year
since 2009
Carbon Aerogels 3
Carbon Nanotube-Based Aerogels and Xerogels
The rst report of a CNT-based aerogel was published by Bryning et al. in 2007
(2007). Their synthesis involved freeze-drying an aqueous CNT suspension, which
yielded a highly conductive, yet fragile network structure that was stabilized only by
van der Waals forces between neighboring CNTs. With the addition of a polymer
binder, the mechanical strength of the aerogel increased signicantly, but electrical
conductivity exhibited a precipitous decline. Since that initial report, efforts in this
area have sought to create CNT architectures that simultaneously exhibit good
transport and mechanical properties, along with tunable surface area and density.
Carbon Nanotube-Based Aerogels
One of the critical aspects in the design of CNT-based aerogels is the selection of the
binder or glue(Morris et al. 1999) used to reinforce these low-density structures.
The utilization of solgel-derived carbon to cross-link the CNTs offers the opportu-
nity to enhance the mechanical properties of the network architecture while
maintaining bulk electrical conductivity. This approach was used to prepare a new
class of ultralow-density (as low as ~10 mg cm
) nanoporous monoliths made of
single-walled carbon nanotubes (SWNT) of macroscopic dimensions with unprec-
edented properties (Worsley et al. 2009a). These SWNT-based carbon aerogels
(SWNT-CA) are the stiffest low-density nanoporous solids reported and exhibit
elastic behavior up to compressive strains as large as ~80 %.
These CNT aerogels were prepared using highly puried SWNT as these tubes
have some acid functionalization allowing them to be readily suspended in water
using sonication without the surfactants or other additives that are typically required
to disperse CNTs in aqueous media. To reinforce the CNT network, organic solgel
chemistry (Pekala 1989) was used as a means to form carbonaceous interconnections
between the CNTs. By introducing low concentrations of the solgel precursors
(resorcinol and formaldehyde) to a suspension of highly puried SWNTs, polymer-
ization occurs primarily on the walls of the CNT bundles and, more importantly, at
the junctions between adjacent bundles to form an organic binder. Thermal treatment
of the resulting aerogel is then used to convert the organic binder to carbon. With this
approach, a series of monolithic SWNT-CAs with CNT content ranging from 0 to
~60 wt% (relative to the carbon binder) were prepared. Interestingly, the densities of
CNT monoliths decreased with increasing CNT content due to smaller volumetric
shrinkage of the parts during both supercritical drying and pyrolysis. Scanning
electron microscopy shows that the network of SWNT-CA is comprised of randomly
interconnected lament-like struts with diameters that range from 5 to 40 nm and
lengths of ~5001000 nm (Fig. 3). Closer examination of these laments by
transmission electron microscopy shows that the CNTs are coated with a thin layer
of carbon, indicating that nucleation and growth of the solgel polymer occurred on
the surfaces of the CNTs.
4 M.A. Worsley and T.F. Baumann
The utilization of the carbonaceous binder affords monolithic CNT aerogels that
simultaneously exhibit exceptional mechanical and electrical properties. In Fig. 4,
the elastic moduli of the SWNT-CA are compared with some other porous carbon
materials, such as traditional RF-derived carbon aerogels. The data illustrates the
unprecedented mechanical properties of the SWNT-CA. Indeed, for a given density,
the SWNT-CAs are the stiffest. In fact, at a density of 100 mg cm
, aerogels with
high CNT loadings (over 16 wt%) are ~12 and ~3 times stiffer than conventional
silica and carbon aerogels, respectively (Pekala et al. 1990; Woignier et al. 1998;
Leventis et al. 2002). These SWNT-CAs are also ~3 times stiffer than the super-
stiffalumina nanofoams whose struts have the morphology of curled nanoleaets
(Kucheyev et al. 2006). The inset in Fig. 4shows a sequence of images taken of an
~30 mg cm
aerogel with a CNT loading of 55 wt% before, during, and after
uniaxial loading up to a maximum strain of ~76 %, showing the superelastic
Fig. 3 SEM images (a,b) of CNT aerogels containing (a) 30 wt% and (b) 55 wt% CNTs and TEM
images (c,d) at different magnications of a foam containing 30 wt% CNT (Reprinted from
reference (Worsley et al. 2009a) with the permission of AIP Publishing)
Carbon Aerogels 5
behavior with complete strain recovery that was observed for SWNT-CAs with
densities below ~50 mg cm
In addition to exceptional mechanical properties, these CNT aerogels also exhibit
high electrical conductivity, as measured by the four-probe method. In Fig. 5,
electrical conductivities, σ, of pristine CAs and SWNT-CA at nanotube loading of
30 wt% of similar densities, ρ, are compared. This double-logarithmic plot reveals a
power-law nature of the σ(ρ) dependence: σ/ρ
, with an exponent nof ~1.55 for all
of the CNT-based aerogels tested. While the exponent nis independent of CNT
loading, Fig. 5clearly shows that the electrical conductivity is ~3 times larger for the
aerogels containing 30 wt% CNTs than for pristine CAs of the same density,
indicating that the effective conductivity of network struts, σ
, increases with
increasing CNT loading. This effect is better illustrated in the inset in Fig. 5, showing
that σ
(dened as σ
, where ρ
is the density of aerogel
struts) dramatically increases for CNT loadings above 16 wt%. For an aerogel
containing ~60 wt% CNTs, a vefold increase in conductivity is observed relative
to a CA of equivalent density. This close-to-linear dependence of the effective σ
on CNT loading is expected for struts made of CNT bundles decorated and
Fig. 4 Dependence of Youngs modulus on monolith density for CNT aerogels with different
nanotube loadings. Data for carbon, silica, and alumina aerogels are shown for comparison. The
inset shows the sequence of uniaxial compression of a monolith (30 mg cm
and 55 wt% CNT
loading), illustrating the superelastic behavior with complete strain recovery after compression to
strains as large as ~76% (Reprinted from reference (Worsley et al. 2009a) with the permission of
AIP Publishing)
6 M.A. Worsley and T.F. Baumann
interconnected with carbon nanoparticles. In this case, strut conductivity is domi-
nated by the resistivity of tube bundles and the array of carbonaceous particles
connected in parallel.
Using this surfactant-free approach, lightweight SWNT-CAs with CNT content
greater than 50 wt% and monolithic densities as low as 10 mg cm
can be prepared.
These SWNT-CAs simultaneously exhibited increased stiffness relative to other
porous materials of equivalent density and high electrical conductivity even at low
densities (Worsley et al. 2009a). However, for some applications, even higher
moduli and/or conductivities are desired.
Carbon Nanotube-Based Xerogels
As seen in Figs. 4and 5, properties such as stiffness and conductivity are highly
dependent on the density of the CNT aerogel. Therefore to exceed the performance
of the SWNT-CA, higher density SWNT-based carbon xerogels (SWNT-CX) were
developed (Worsley et al. 2013). Like the SWNT-CAs, graphitic carbon aerogel
particles are used to cross-link the CNT bundles that constitute the 3D network.
However, instead of supercritical drying the initial CNT wet gel to minimize
shrinkage and achieve an ultralow-density part, the wet CNT gel is dried under
ambient conditions to encourage shrinkage due to capillary forces and obtain a
Fig. 5 The dependence of electrical conductivity on the monolith density for CNT aerogels (30 wt%
CNTs) and pristine carbon aerogels. The inset shows a dramatic increase in the effective conductivity
of foam struts σ
with increasing nanotube loading (Reprinted from reference (Worsley
et al. 2009a) with the permission of AIP Publishing)
Carbon Aerogels 7
higher density gel. The ambient drying before carbonization allows the SWNT-CX
to achieve an order of magnitude increase in density over the SWNT-CA while
maintaining the high surface area. As expected, the SWNT-CX exhibits several
orders of magnitude higher electrical conductivity and elastic modulus than those
for the ultralow-density SWNT-CA.
Figure 6shows SEM images of SWNT-CA and SWNT-CX, revealing that both
materials have a similar network of randomly interconnected ligament structures,
consistent with previous SEM studies of the SWNT-CA (Worsley et al. 2009a). A
key qualitative difference between the SWNT-CA and SWNT-CX appears to be the
size of the pores formed by the network of CNT bundles. In particular, the SWNT-CA
(Fig. 6a) appears to contain many pores in excess of 100 nm, while the SWNT-CX
(Fig. 6b) possesses the majority of pores in the 10100 nm range. A more quantitative
analysis of effective pore sizes is provided by nitrogen porosimetry data (Fig. 6c,d).
The nitrogen isotherms (Fig. 6c) indicate that the SWNT-CX does have a larger
mesopore (pores <300 nm) volume (0.9 cm
/g) than the SWNT-CA (0.6 cm
This is also conrmed by pore size distributions shown in Fig. 6d, illustrating that the
bulk of the surface area for the SWNT-CA lies below 10 nm, while the SWNT-CX has
most of its surface area in the 10100 nm range. The pore size distribution of the
SWNT-CX is likely the result of pore collapse within the gel during evaporative drying
under ambient conditions. The densication is signicant, converting a ~30 mg/cm
SWNT-CA, which is over 98 % macroporous, to a ~450 mg/cm
SWNT-CX, which is
only 50 % macroporous. The surface area of the SWNT-CX is maintained at ~590 m
Structural integrity and electrical conductivity are important considerations in
designing porous structures for use as electrodes. Therefore, we have evaluated both
mechanical and electrical properties of the SWNT-CX. Figure 7shows that the
SWNT-CX, like the SWNT-CA, possesses electrical conductivity about three
times larger than that of the conventional CA without CNTs.
The mechanical deformation behavior of porous solids is strongly dependent on
the monolith density (Worsley et al. 2009a). The SWNT-CX has a Youngs modulus
of ~1.2 GPa and an effective failure stress of ~100 MPa. As mentioned earlier,
Youngs modulus of conventional CAs and SWNT-CAs depends superlinearly on
monolith density with an exponent of ~2.7.
Fig. 7shows that the modulus of the
SWNT-CX scales well with the density as expected for CNT-based materials. This
suggests that ligament connectivity does not change during ambient-drying-induced
densication of SWNT-CA to SWNT-CX.
Carbon Nanotube-Based Aerogel Composites
One promising application of CNT aerogels is in the area of conductive composite
materials. Traditionally, effort has been focused on various methods to disperse
CNTs into the target matrix, with limited success. Effective dispersion of CNTs is
extremely challenging and many times requires specic functionalization of the
CNTs, which can be expensive, time-consuming, and is not universally applicable.
On the other hand, the ultralight and robust CNT aerogels can serve as scaffolds for
the preparation of novel CNT composites. As the CNT network is already
8 M.A. Worsley and T.F. Baumann
established, the aerogel can simply be impregnated through the wicking process
(Wang et al. 1993) with the matrix of choice, ranging from inorganic sols to polymer
melts to ceramic pastes, to prepare a variety of conductive CNT composites (Worsley
et al. 2009b,2011a). As a rst example, the SWNT-CA is used as a scaffold for the
synthesis of a stiff, highly conductive poly(dimethylsiloxane) (PDMS) composite.
This polymer/CNT composite exhibits ~300 % increase in the elastic modulus
relative to the unloaded PDMS elastomer and electrical conductivity over
1 Scm
, the highest conductivity reported for a polymer/SWNT composite at this
CNT loading level (1.2 wt% or 1 vol.%) (Winey et al. 2007; Mathur et al. 2008).
CNT-polymer composites were prepared by immersing the as-prepared SWNT-
CA in the polymer resin, Dow Corning Sylgard 184, prior to cure. The immersed
SWNT-CA was placed under vacuum until no more air escaped from the scaffold,
Fig. 6 SEM of (a) SWNT-CA and (b) SWNT-CX. (c) Nitrogen adsorption/desorption isotherms
and (d) pore size distributions for SWNT-CA and SWNT-CX (Reproduced by permission of The
Electrochemical Society)
Table 1 Physical properties of SWNT-CA and SWNT-CX (Reproduced by permission of The
Electrochemical Society)
Surface area,
conductivity, S/cm
modulus, MPa
SWNT-CA 0.03 589 1.12 1.1
SWNT-CX 0.45 593 67 1200
Carbon Aerogels 9
suggesting full inltration of the resin. The inltrated SWNT-CA was then cured at
60 C to produce the composite. The dimensions of the composite were approxi-
mately equal to those of the initial SWNT-CA.
SEM images of PDMS/SWNT-CA composites show that the SWNTs are
homogenously distributed throughout the polymer matrix, suggesting that there is
good wetting at the PDMS/SWNT-CA interface and that the CNT-based scaffold is
intact after inltration and curing (Fig. 8). This observation is supported by the fact
that the electrical conductivity of the SWNT-CA scaffold is maintained even in a
fully dense insulating matrix (Table 2). The conductivity of these polymer compos-
ites (1 Scm
) represents the highest conductivity reported for a polymer/SWNT
composite prepared at such a low CNT loading level (1.2 wt% or 1 vol.%) (Winey
et al. 2007; Mathur et al. 2008). Interestingly, the electrical conductivity of this
composite is on par with the highest reported value for a polymer/MWNT at a similar
~1 wt% MWNT loading (Grossiord et al. 2008). As SWNTs typically contain some
fraction of semiconducting tubes, as compared to MWNTs, which presumably are all
metallic, one might expect a higher conductivity in the MWNT composite with
similar CNT loadings. This observation highlights the need for further study in this
area and suggests that even larger improvements in the conductivity of polymer
composites are possible.
Nanoindentation measurements show that the PDMS/SWNT-CA exhibits elastic
behavior with an ~300 % increase in Youngs modulus as compared to the case of
PDMS (Fig. 9). The observed enhancement in modulus is consistent with the
increase expected based on the Halpin-Tsai model for a nanotube bundle aspect
ratio of ~100 (Halpin 1969). Similar increase in modulus was observed by Dyke
et al. for a PDMS/SWNT composite prepared with 1 wt% loading of surface-
functionalized SWNT (Dyke and Tour 2004). The improved modulus is also con-
sistent with the observation of a polymer shell that coats the CNT bundles in the
SEM images (Fig. 8). The presence of the polymer shell suggests strong bonding at
Fig. 7 Log-log plots comparing the dependence of (a) electrical conductivity and (b) Youngs
modulus on density for carbon aerogels, CNT aerogels, and CNT xerogels (Reproduced by
permission of The Electrochemical Society)
10 M.A. Worsley and T.F. Baumann
the PDMS/SWNT-CA interface, a key element in successful reinforcement
(Thostenson et al. 2001). These results highlight the effectiveness of using a
premade CNT scaffold for structural reinforcement.
To highlight the universal applicability of using CNT aerogels for conductive
composites, the next example describes the fabrication of novel oxide/SWNT-CA
Fig. 8 SEM images (under different magnications) of conductive PDMS/SWNT-CA composites
(Reproduced from reference (Worsley et al. 2009b) with the permission of the Royal Society of
Table 2 Physical properties of SWNT-CA, polymer, and polymer-CNT composite (Reproduced
from reference (Worsley et al. 2009b) with the permission of the Royal Society of Chemistry)
Material CNT, vol% (wt%) Density, g/cm
E, MPa σ, Scm
SWNT-CA 1 (55) 0.03 1.1 1.12
PDMS 0 (0) 1.04 4.2 <0.001
PDMS/SWNT-CA 1 (1.3) 1.01 14 1.00
Fig. 9 Partial
loaddisplacement curves for
PDMS with and without CNT
aerogel. The inset shows
depth proles of the
indentation modulus
(Reproduced from reference
(Worsley et al. 2009b) with
the permission of the Royal
Society of Chemistry)
Carbon Aerogels 11
composites through the solgel deposition of oxides on the surface of nanoligaments
of SWNT-CA. Unlike the nonporous polymer/SWNT-CA composites, the oxide is
deposited as a conformal overlayer on the primary ligament structure of the SWNT-
CA support, thus preserving open porosity within the monolithic part. This work
specically focuses on nanoporous composites prepared with SiO
coatings due to the technological interest in each of these materials. In particular,
composites of SiO
and CNTs (Ding et al. 2009;Shinetal.2007; Hernadi et al. 2003;
Fu et al. 2002) possess desirable properties for a wide range of applications such as
electrochemical devices (Gavalas et al. 2001), catalysis (Lee et al. 2008), sensors (Guo
et al. 2009), optoelectronics (Olek et al. 2006), and separations (Zhang et al. 2010).
/CNT composites (Han and Zettl 2003) have shown potential as sensors (Yang
et al. 2010;Gongetal.2008), catalyst supports (Hsu et al. 2010;Duetal.2009), and
Li-ion battery electrodes (Du et al. 2010; Zhu et al. 2010a;Fuetal.2009; Chen
et al. 2008a,b;Anetal.2007; Xie and Varadan 2005), while TiO
/CNT composites
(Shin et al. 2007; Hernadi et al. 2003; Liu and Zeng 2008;Yuetal.2007)areof
interest due to their potential impact in elds of catalysis (An et al. 2007; Wang
et al. 2005,2008a; Orlanducci et al. 2006), photoelectronics (Yang et al. 2007), and
energy storage (Wang et al. 2008b;Mishraetal.2008). Structural characterization of
the oxide/SWNT-CA composites shows that, in each case, solgel deposition of the
oxide coating occurs primarily on the surfaces of CNT ligaments throughout the
aerogel monolith. In addition, these composite structures retain the high electrical
conductivity of the SWNT-CA support and exhibit signicant enhancements in
mechanical properties relative to the uncoated aerogel structure. The approach
described here provides a straightforward route to the design of porous and monolithic
oxide/CNT composites for use in a variety of applications.
The oxide/SWNT-CA composites were prepared through deposition of an oxide
coating over the inner surface area of the SWNT-CA framework using solgel
chemistry. The oxide solgel solutions were synthesized using literature methods
(Kucheyev et al. 2005; Baumann et al. 2005; Iler 1979). The SiO
solgel was
prepared via traditional one-step base-catalyzed alkoxide solgel chemistry using
tetramethoxysilane (4.1 g), water (1.5 g), ammonium hydroxide (30 %, 200 ml), and
methanol (24 g) (Iler 1979). The TiO
solgel was prepared via a two-step process
involving acid-catalyzed hydrolysis of titanium (IV) ethoxide (1 g) using water
(85.7 ml), hydrochloric acid (37 %, 71.4 ml), and ethanol (3.57 g), followed by
base-initiated gelation using propylene oxide (0.357 g) (Kucheyev et al. 2005). The
solgel was prepared via an epoxide-initiated gelation method using tin
chloride pentahydrate (0.56 g), trimethylene oxide (1.03 g), ethanol (7 g), and
water (5 g) (Baumann et al. 2005). Composites were synthesized by inltration of
SWNT-CA monoliths by the oxide solgel solutions prior to gelation. The SWNT-
CAs were immersed in the solgel solutions and placed under vacuum until no more
air escaped from the scaffolds, indicating full penetration of the sol. The concentra-
tion of inorganic precursors was kept low to promote the growth of the condensed
inorganic phase primarily on the surfaces of the SWNT-CA framework while
minimizing gelation in the free pore volume of the aerogel. The inltrated SWNT-
CAs were then cured at room temperature for 72 h to produce the wet oxide/SWNT-
12 M.A. Worsley and T.F. Baumann
CA gels. The wet oxide/SWNT-CA gels were dried using supercritical extraction
with liquid CO
to yield the nal dry oxide/SWNT-CA composites.
Examination of the composite structures by SEM (Fig. 10) shows that the porous
network structure of the SWNT-CA support is maintained after the oxide deposition
and drying process. Additionally, in each case, the deposited oxide appears to form a
uniform coating on the surfaces of the CNT ligaments. For example, the SEM images
for the SiO
/SWNT-CA composite show that the deposited SiO
aerogel particles
preferentially coat the CNT bundles, as very few unsupported SiO
particles were
observed (Fig. 10a,b). Similarly, the SEM images for the SnO
/SWNT-CA show that
the majority of SnO
particles are associated with the CNT network (Fig. 10c,d).
Unlike the SiO
composite, however, the deposited SnO
particles are crystalline
with diameters of ~35 nm, consistent with previous reports for SnO
prepared using a similar solgel formulation (Baumann et al. 2005). The TiO
particles, as deposited, are amorphous but can be converted to the anatase phase
after calcination at 320 C for 5 h in air. After annealing, the TiO
nanocrystals are
easily distinguishable by SEM from the underlying support structure (Fig. 10e,f).
Not surprisingly, the gas adsorption properties of the oxide/SWNT-CA compos-
ites are quite different from those of the uncoated support material. As shown in
Fig. 11, the nitrogen isotherms for these composite structures are consistent with
materials comprised of both meso- and macropores. Since the SWNT-CA support is
primarily a macroporous structure, the mesoporosity in these materials can be
attributed to the porous oxide coating. Each of the composite structures exhibits
larger surface area and mesopore volume than those respective values for the SWNT-
CA support, which are similar to surface areas observed in pure metal oxide aerogels
(Kucheyev et al. 2005; Baumann et al. 2005; Iler 1979; Pajonk 1997; Taguchi and
Schuth 2005). Notably, the BET surface area (742 m
/g) and mesopore volume
(2.2 cm
/g) measured for the SiO
/SWNT-CA composite are considerably larger
than the values measured for the SWNT-CA support (162 m
/g, 0.3 cm
/g) and on
par with values observed in the pure SiO
aerogel (670 m
/g, 1.8 cm
/g). This
change in textural properties has been observed with other SiO
/CNT composites
(Zhang et al. 2010). The SnO
composite also exhibits signicant increases in BET
surface area (349 m
/g) and mesopore volume (0.9 cm
/g) relative to the support
structure. Interestingly, the crystalline TiO
/SWNT-CA composite has a greater
mesopore volume than the amorphous composite (1.1 vs. 0.6 cm
/g), while the
BET surface areas remain similar (197 and 202 m
/g for crystalline and amorphous
composites, respectively). Thus, crystallization of the deposited TiO
particles was
achieved without loss of accessible surface area. The high surface areas and large
mesopore volumes associated with these composite materials would be expected to
be benecial for energy storage, catalyst, and sensing applications.
Structural integrity and electrical conductivity are also important considerations
in designing porous structures for use as catalyst supports and electrodes. Therefore,
the mechanical and electrical properties of the composite structures were evaluated
(Table 1and 2and Fig. 8). The electrical conductivities of each of the oxide-coated
composites are similar to that of the uncoated SWNT-CA structure, suggesting that
the conductive CNT network is intact after deposition and supercritical drying. This
Carbon Aerogels 13
observation is consistent with our previous results for polymer/SWNT-CA compos-
ites (Worsley et al. 2009b).
In any analysis of the mechanical deformation behavior of porous solids, the rst
parameter to consider is the monolith density since mechanical properties depend
superlinearly on the density (Worsley et al. 2009a). Due to differences in the
geometry and connectivity of ligaments, different aerogels exhibit different scaling
behavior of E on the monolith density (ρ). In particular, for TMOS-derived base-
Fig. 10 SEM images of SiO
/SWNT-CA (a,b), SnO
/SWNT-CA (c,d), and TiO
/SWNT-CA (e,
f) at low and high magnications (Reprinted with permission from reference (Worsley et al. 2011a).
Copyright 2016 American Chemical Society)
14 M.A. Worsley and T.F. Baumann
catalyzed SiO
aerogels, as studied in this work, the E scales as E ~ ρ
et al. 1998), while for our CNT-based and conventional carbon aerogels, E ~ ρ
(Worsley et al. 2009a). The SiO
/SWNT-CA composite aerogel with a monolith
density of 80 mg/cm
is made of 30 mg/cm
of carbon and 50 mg/cm
of SiO
. The
analysis for SiO
/CNT composites is further simplied by the fact that the mass
densities of ligaments in both aerogels are similar (~2 g/cm
). If such a composite
aerogel were made from two interpenetrating and not cross-linked networks of
carbon and SiO
nanoligaments with monolith densities of 30 and 50 mg/cm
, the
contribution to the elastic modulus from carbon and SiO
networks would be ~1 and
0.1 MPa, respectively (Worsley et al. 2009a; Woignier et al. 1998). Hence, a much
larger Youngs modulus of ~7 MPa of the SiO
/SWNT-CA composite (Table 1)is
consistent with electron microscopy observations (Fig. 10) that the composite
aerogel has morphology of the carbon scaffold with SiO
particles decorating carbon
ligaments rather than of two interpenetrating but poorly cross-linked SiO
carbon networks.
This scaling law analysis could be extended further by noting that a CNT-based
aerogel with a density of 80 mg/cm
has a Youngs modulus of ~15 MPa (Worsley
et al. 2009a), which is close to the modulus of the SiO
/CNT composite (~7 MPa). A
factor of two lower modulus of the composite could be attributed to signicantly
weaker atomic bonds in SiO
(and, hence, a lower E) as compared to the case of
graphitic carbon. Indeed, the Youngs modulus of full-density amorphous silica is
~70 GPa as compared to typical ~200400 GPa modulus values for graphitic carbon
bers. Hence, coating of carbon aerogel ligaments with SiO
nanoparticles has a
mechanical reinforcement effect comparable to that of increasing the density of the
carbon scaffold, suggesting good adhesion between the deposited SiO
and the SWNT-CA support.
An even greater increase in the elastic modulus is observed in TiO
composites (Table 3). Since the density of full-density TiO
is about twice that of
Fig. 11 (a) Nitrogen adsorption/desorption isotherms and (b) pore size distributions for SWNT-
CA and SiO
/SWNT-CA (Reprinted with permission from reference (Worsley et al. 2011a). Copy-
right 2016 American Chemical Society)
Carbon Aerogels 15
graphitic carbon and Youngs moduli of full-density TiO
and graphitic carbon are
similar, a uniform coating of a SWNT-CA scaffold with an initial density of 30 mg/
with TiO
with a volumetric weight density of 50 mg/cm
is comparable to a
case of a purely carbon nanofoam with a monolithic density of ~30 + 50*2/
4=55 mg/cm
. According to the E scaling behavior for CNT aerogels (Worsley
et al. 2009a), such a carbon aerogel has E ~ 6 MPa, which is three times lower than
E for the TiO
/SWNT-CA composite (Table 1). This could be attributed to additional
cross-linking of ligaments by TiO
particles as opposed to a less efcient process of a
uniform coating of ligaments of the carbon scaffold. The cross-linking effectively
changes the connectivity of the ligament network and improves mechanical proper-
ties and their scaling behavior.
Table 3shows that, in contrast to the cases for SiO
and TiO
, the mechanical
properties of the SnO
/SWNT-CA composite are dominated by those of the CNT
scaffold. This difference in reinforcement may be related to different amounts of
oxide deposited on the SWNT-CA (33 wt% for SnO
vs. 70 wt% for SiO
and TiO
resulting in an increase of the monolith density by ~17 at.% for SnO
versus ~167
and ~83 at.% for SiO
and TiO
, respectively. Alternatively, such differences could
also arise from different interfacial interactions between the deposited oxide and the
SWNT-CA support due to the different reaction chemistries used to deposit oxide
coatings. Additional studies of deformation behavior of carbon aerogels with vari-
able loading of SnO
could provide insight into the structureproperty correlation of
/SWNT-CA composites.
Composites of TiO
and carbon have been shown to exhibit higher photocatalytic
activity over a wider absorption band relative to TiO
by itself. Therefore, we have
also measured the UVvis absorption spectrum of the annealed TiO
composite as a suspension in ethanol. As shown in Fig. 12, the spectrum for the
/SWNT-CA material shows not only the well-known TiO
absorption band
around 260 nm but also absorption far into the visible at 500 nm when the
contribution from the CNT scaffold is removed (the inset in Fig. 12). This widening
of the absorption band in TiO
/C composites is well documented in the literature and
is another reason why TiO
/CNT composites are of such high interest (Sakthivel and
Kisch 2003).
Table 3 Physical properties of SWNT-CA scaffold, bulk SiO
, SiO
/SWNT-CA, bulk SnO
/SWNT-CA, bulk TiO
, and TiO
/SWNT-CA composites (Reprinted with permission from
reference (Worsley et al. 2011a). Copyright 2016 American Chemical Society)
Material CNT, vol% (wt%) Density, g/cm
E, MPa σ, Scm
SWNT-CA 1 (55) 0.028 1.2 0.1 1.12
0 0.120 2.4 0.2 <0.001
/SWNT-CA 1 (16) 0.080 7.3 0.4 1.00
0 0.204 3.6 1.7 <0.001
/SWNT-CA 1 (33) 0.046 1.0 0.2 1.00
0 0.193 3.5 0.2 <0.001
/SWNT-CA 1 (16) 0.082 17.3 1.7 1.00
16 M.A. Worsley and T.F. Baumann
In summary, novel oxide/CNT nanocomposites were fabricated through the
deposition of oxide coatings on monolithic SWNT-CAs supports. Inltration and
deposition of the oxides were achieved with little to no degradation of the extended
CNT network of the support, yielding highly conductive nanocomposites. Signi-
cant mechanical reinforcement was also observed. In addition, the oxide/CNT
nanocomposites exhibited high surface areas and large internal pore volumes. The
monolithic nature of these oxide/SWNT-CA assemblies should prove advantageous
for a number of applications, such as battery electrodes, sensing devices, and
catalysts. These results, along with those for the polymer/SWNT-CA composites,
show the versatility and potential of the CNT aerogel scaffold approach.
Graphene-Based Aerogels and Xerogels
Given the relatively cheap precursors and mild synthesis conditions, GO-based wet
chemistry methods are easily the most popular routes for producing graphene
aerogels (Zhu et al. 2010b; Chen et al. 2012). Unlike CNTs, GO is readily
dissolved/suspended in aqueous solutions making it relatively easy to work with
and more environmentally friendly. However, due to the high degree of oxidation
(i.e., sp
carbon) present in GO compared to CNTs, some reduction step, typically
chemical or thermal, is required to recover the desired graphene-like properties.
Here, we present some common solgel methods to produce graphene-based
aerogels (Worsley et al. 2010,2011b,2012) and xerogels (Worsley et al. 2014),
along with characterization and their use in some well-known applications.
Fig. 12 UVvis absorption
spectra for TiO
and SWNT-CA. Spectra are
offset for clarity. Inset shows
contribution of TiO
nanoparticles in TiO
CA (SWNT-CA spectrum
subtracted) (Reprinted with
permission from reference
(Worsley et al. 2011a).
Copyright 2016 American
Chemical Society)
Carbon Aerogels 17
Graphene Aerogels via Resorcinol-Formaldehyde Solgel Chemistry
In 2010, the rst synthesis of a graphene aerogel that used covalent carbon cross-
links between graphene sheets instead of physical bonds was reported (Worsley
et al. 2010). The fabrication scheme involved the gelation of a GO suspension via RF
solgel chemistry (Al-Muhtaseb and Ritter 2003). The GO was produced by Hum-
mers method (Hummers and Offman 1958) and the suspension was prepared by
ultrasonication. The molar ratio of R:F was 1:2, the reactant concentration in the
starting mixture was 4 wt% RF solids, and the concentration of GO in suspension
was 1 wt%. The molar ratio of reactant:catalyst (sodium carbonate) was 200:1. The
solgel mixture was cured in sealed glass vials at 85 C. After gelation, the wet
GO-RF gels were removed from the glass vials and washed in acetone to remove
water from the pores. Supercritical CO
was used to dry the GO-RF gels, and
pyrolysis at 1050 C under nitrogen yielded the nal graphene aerogel. Energy
dispersive X-ray analysis conrmed the successful reduction of the GO-RF gel
showing a drop in atomic oxygen from 17 % to 1 %.
SEM images of the graphene aerogel show a 3D network of randomly oriented
sheetlike structures (Fig. 13ab) similar to those seen in previous reports of ther-
mally reduced GO (Yoo et al. 2008). The lateral dimensions of the sheets ranged
from hundreds of nanometers to several microns. Within the assembly, the sheets are
thin enough to be transparent to the electron beam. Transmission electron micro-
graphs (TEM) reveal a wrinkled paper-like texture to the sheets (Fig. 13c,d), again
consistent with previous reports (McAllister et al. 2007). It is important to note that,
in both the SEM and TEM images, we do not observe carbon nanoparticles from the
RF polymer decorating the surfaces of the graphene sheets, despite the fact that over
half of the weight in the reduced GO-RF structure can be attributed to carbon from
the RF polymer (56 wt% from RF vs. 44 wt% from GO). This is in sharp contrast to
what occurs in materials prepared at higher RF:GO ratios, or in CNT aerogels, where
the carbonized RF is clearly distinguishable (Worsley et al. 2009c). This observation
suggests that the carbon junctions are effectively incorporated into the extended
graphene framework. When added at sufciently low concentrations to the GO
suspension, polymerization of the organic precursors likely occurs preferentially at
the oxygen functionalities of the GO to form covalent interconnections between
individual sheets. Simultaneous carbonization of the RF junctions and thermal
reduction of the GO apparently blends these two materials into a single structure,
yielding the graphene aerogel, which exhibited large surface areas (>500 m
/g) and
electrical conductivities orders of magnitude higher than those of graphene aerogels
made with physical cross-links alone (Wang and Ellsworth 2009).
While the RF likely plays a signicant role in cross-linking the GO when present,
GO has a number of functional groups (e.g., epoxide, hydroxyl, carboxyl) that could
result in cross-linking without the RF under the appropriate conditions. A better
understanding of how the cross-linking agent impacts the nal aerogel properties
represents another signicant development in the graphene aerogel research. A
recent study with varied RF content (04 wt%) shows how the synthetic parameters
18 M.A. Worsley and T.F. Baumann
affect the microstructure, the nature of the carbon cross-links, and nally the bulk
properties (e.g., surface area, electrical conductivity) of the graphene aerogel.
Gelation of the GO suspension occurred across the full range of RF content
(04 wt%). This result supports the hypothesis that cross-linking can occur not only
via RF solgel chemistry but also solely between the functional groups present on
suspended GO sheets. Previous work also reports the gelation of GO suspensions,
but has been limited to methods requiring high temperature and pressure (e.g., in an
autoclave) (Tang et al. 2010;Xuetal.2010), or producing thin lms (1 μm or less)
(Liu and Seo 2010). In contrast, the present base-catalyzed method was achieved at
less than 100 C in glass vials, and could conceivably be accomplished at room
temperature. Further investigations of the low-temperature gelation of GO suspen-
sions are in progress and will be presented in a subsequent report. After gelation, all
of the samples were dried using supercritical CO
and pyrolyzed at 1050 C under
to reduce the GO and RF to carbon.
Details of the microstructure of the graphene-based aerogels are revealed by SEM
(Fig. 14). The micrographs show a signicant change in morphology as the RF
content is reduced. All of the graphene-based aerogels exhibit a sheetlike structure
Fig. 13 SEM of the graphene aerogel at low (a) and high (b) magnication. TEM of the graphene
aerogel at low (c) and high (d) magnication. Black arrow denotes holey carbon on TEM grid
(Reprinted with permission from reference (Worsley et al. 2011b). Copyright 2016 American
Chemical Society)
Carbon Aerogels 19
similar to that observed in other graphene assemblies (Wang and Ellsworth 2009;
Tang et al. 2010; Xu et al. 2010). In addition, no carbon nanoparticles from the RF
polymer were observed on the surfaces of the graphene sheets up to 4 wt% RF
content. However, as the RF content decreases, the features of the aerogel appear to
become much ner. Specically, the sheets in the aerogel appear more transparent
(e.g., thinner) and the size of the voids between sheets appears smaller, particularly
in the aerogel without any resorcinol. This observation would be consistent with
maintaining a higher degree of exfoliation in the nal graphene assembly when a
lower RF content is used in the gelation process.
X-ray powder diffraction (XRD) spectra for the graphene-based aerogels with
(4 wt%) and without the RF are shown in Fig. 15. The diffraction pattern for GO
shows a characteristic peak at ~12. This feature is not present in either of the
graphene-based aerogels. The material with 4 wt% RF content shows a broad peak
from 12 to 20and a sharper feature at ~28. The broad peak is indicative of
disordered few-layer graphene sheets which retain the GO interlayer separation
(McAllister et al. 2007). The 100/101 graphite peaks at ~28(interlayer separation
of graphite) are attributable to the heat-treated carbon derived from a disordered or
organic precursor (i.e., RF polymer) (Takai et al. 2003). The sample without RF
Fig. 14 SEM images of the graphene aerogels with (a) 4 wt%, (b) 2 wt%, and (c) 0 wt% initial RF
content (Reprinted with permission from reference (Worsley et al. 2011b). Copyright 2016 Amer-
ican Chemical Society)
20 M.A. Worsley and T.F. Baumann
shows no signicant peaks indicating an amorphous structure. This is similar to what
is observed with graphene sheets that are very well exfoliated (McAllister
et al. 2007). Though some layering is likely still present, it is disordered enough
not to produce any diffraction peaks. Therefore, the XRD results also suggest
superior exfoliation of the graphene sheets is maintained in the aerogel when the
RF is omitted.
Figure 16 displays C(1 s) X-ray absorption spectroscopy (XAS) spectra recorded
at a 45angle of incidence for graphene-based aerogels of different RF content and a
freshly cleaved HOPG standard. The sharp resonance observed at ~285.4 eV in all of
the spectra arises from the C(1 s) !π
transitions anticipated for materials
containing predominantly sp
carbon (Stöhr 1992). Meanwhile, the resonance at
~291.5 eV is attributed to a core-hole exciton state (Ma et al. 1993) and the series of
broad resonances that follow at higher energies arise primarily from C(1 s) !CC
transitions. Neither the graphene aerogel prepared in the absence of RF nor the
HOPG standard exhibits appreciable resonances in the characteristic range for
spectral features associated with CHσ
-transitions, ~287290 eV (Stöhr 1992).
As such, the proportion of carbon atoms bonded to hydrogen (i.e., those at the edge
of a domain) in these materials resides close to or below the detection limits of the
XAS measurement (~2 %), and the domain/crystallite size must be large enough
that, on average, carbon atoms bonded only to carbon (i.e., at the interior of a
domain) predominate by at least 98 %. In contrast, the inclusion of 4 % wt RF in
the graphene-based aerogel leads to a broad absorption onset at lower energy than
the excitonic feature, beginning at ~289.4 eV, which indicates an increase in the
Fig. 15 XRD patterns of the graphene aerogels with 0 and 4 wt% initial RF content. Graphene
oxide is included for reference. Patterns are offset for clarity (Reprinted with permission from
reference (Worsley et al. 2011b). Copyright 2016 American Chemical Society)
Carbon Aerogels 21
proportion of carbon bonded to hydrogen versus the HOPG standard and graphene
aerogel prepared without RF. This observation in such a highly cross-linked network
(Fig. 1) would support the hypothesis that the cross-links in the graphene aerogel
prepared without RF are predominantly sp
in nature. With all of the junctions
present in this aerogel, signicant sp
cross-linking would be apparent in the XAS
spectra, as is the case for the aerogel prepared with RF. We propose that, much like
the CH
-cross-links between individual aromatic rings in RF gels are broken during
pyrolysis allowing the formation of aromatic multiring structures (e.g., sp
on the nanoscale (Al-Muhtaseb and Ritter 2003; Kuhn et al. 1998), similar organic
cross-links are likely formed between the GO sheets during gelation which can then
be broken during pyrolysis to allow sp
bonding between the reduced GO sheets.
Analysis of the resonance associated with the core-hole exciton provides addi-
tional insight into the composition and structure of the graphene-based aerogels. The
intensity of the excitonic resonance in the XAS spectra of the two aerogels is greatly
reduced with respect to the HOPG. We postulate that this difference arises due to a
more heterogeneous chemical and physical structure in the aerogels than the HOPG
standard. Comparison between the two aerogel spectra indicates that the excitonic
resonance has a lower intensity and is less well resolved for the sample containing
RF, which we accordingly attribute to the presence of greater degree of
Nitrogen porosimetry data indicate that the RF content plays a critical role in the
textural properties of the graphene-based aerogels (Fig. 17a). Type IV adsorption/
desorption isotherms were observed for all the graphene-based aerogels, indicating
that signicant mesoporosity exists. Consistent with the sheetlike structure observed
in the electron micrographs, a Type 3 hysteresis loop (IUPAC classication) at high
relative pressure points to adsorption within aggregates of platelike particles for this
material. As the RF content is reduced, total adsorption, as well as the size of the
Fig. 16 XAS spectra
recorded at the carbon K-edge
for graphene aerogels with
0 and 4 wt% initial RF content
and a freshly cleaved HOPG
reference (Reprinted with
permission from reference
(Worsley et al. 2011b).
Copyright 2016 American
Chemical Society)
22 M.A. Worsley and T.F. Baumann
hysteresis loop, grows. These phenomena translate into signicantly increased BET
surface area (1199 m
/g) and pore volume (6.4 cm
/g) in the aerogel without
resorcinol (Table 4). The large increase in surface area is likely due to the higher
degree of exfoliation in the aerogel without RF content. While the theoretical limit
for an individual graphene sheet (>2500 m
/g) (Peigney et al. 2001) was not
achieved, surface areas approaching this value appear to be in reach by improving
the exfoliation of graphene sheets in the assembly. The pore size distribution
(Fig. 17b) shows that the peak pore diameter drops from 13 nm at 4 wt% RF content
to 7 nm at 0 wt% RF content. The shift in pore size distribution with decreasing RF
content is consistent with the change in feature sizes observed by SEM for these
The electrical conductivities of the graphene-based monoliths are also shown in
Table 4for the range of RF contents. A modest decrease in conductivity was
observed with the decrease in RF content. This is likely due to the lower density
of the pure graphene aerogel, given the strong dependence of electrical conductivity
on the bulk density of porous carbons.
That said, the conductivity of the pure
graphene aerogel is still comparable to the aerogels cross-linked with RF-derived
carbon and is orders of magnitude higher than previously reported graphene assem-
blies (Tang et al. 2010; Xu et al. 2010).
In summary, graphene aerogels were successfully fabricated with a range of RF
content. In particular, low-temperature gelation of the GO suspension was realized
without resorcinol yielding a graphene aerogel free of RF-derived carbon. The
decreasing RF content produced a graphene assembly with a higher degree of
exfoliation and less CH bonding than observed at higher RF content. The resulting
aerogel exhibited extraordinarily high surface area (~1200 m
/g) and large pore
volume (~6 cm
/g) while maintaining the high conductivity observed in the
RF-derived graphene aerogels. Given these novel properties, these graphene
Fig. 17 (a) Nitrogen adsorption/desorption isotherms and (b) normalized pore size distribution
plots for the graphene aerogels (Reprinted with permission from reference (Worsley et al. 2011b).
Copyright 2016 American Chemical Society)
Carbon Aerogels 23
aerogels should be attractive candidates for energy storage, sensing, and catalytic
Graphene Aerogels from Pure Graphene Oxide Suspensions
As shown in the previous section, under the right conditions, the gelation of an
aqueous GO suspension can proceed without the aid of traditional RF solgel
chemistry. In fact, reports have shown that GO suspensions can undergo gelation
under a number of different conditions (Bai et al. 2011), including low and high pH
conditions, as well as using high temperature (>150 C) hydrothermal methods
(Worsley et al. 2012; Xu et al. 2010; Bai et al. 2011; Zhang et al. 2011; Niu
et al. 2012; Qiu et al. 2012; Li et al. 2014). A signicant challenge for graphene
aerogel synthesis lies in being able to initiate enough self-assembly and reduction of
the GO sheets to build a mechanically robust, electrically conductive network
without a large degree of sheet restacking so that surface area is not compromised.
A fairly reliable method for achieving this goal is to perform the gelation under basic
conditions which uses the various functional groups (e.g., epoxide, hydroxide)
abundant in GO sheets serve as chemical cross-linking sites (Worsley et al. 2012).
After thermal reduction, these cross-linkers served a dual purpose. These sp
bridges provide the mechanical reinforcement and electrical conductivity, while
simultaneously serving as spacers, which prevent restacking so that large surface
areas can also be achieved.
This high surface area, electrically conductive, and mechanically robust graphene
aerogel was prepared by gelation of a GO suspension under basic conditions. The
aqueous GO suspension (12 wt%) was prepared by ultrasonication. In a glass vial,
3 ml of the GO suspension was mixed with 500 μl concentrated NH
OH. The vial
was sealed and placed in an oven at 85 C overnight. The resulting wet gel was
washed in deionized water to purge NH
OH followed by an exchange of water with
acetone inside the pores. Supercritical CO
was used to dry the gels that were then
converted to the nal graphene aerogel by pyrolysis at 1050 C under nitrogen.
Densities of the black monoliths were 80100 mg/cm
Solid-state nuclear magnetic resonance (NMR) characterization was used to gain
insight into the types of functional groups in GO involved in the cross-linking
process as well as to follow the reduction of GO to graphene (Fig. 18a). The GO
powder contains signicant epoxide and hydroxyl functionality as evidenced by
Table 4 Physical properties of graphene-based aerogels (Reprinted with permission from refer-
ence (Worsley et al. 2011b). Copyright 2016 American Chemical Society)
Initial RF
content, wt%
BET surface
area, m/g
volume, cm
conductivity, S/m
4 0.025 584 2.9 87
2 0.019 762 3.3 55
0 0.016 1199 6.4 25
24 M.A. Worsley and T.F. Baumann
numerous peaks between 50 and 75 ppm, as well as carbonyl groups (168 ppm) and
carbon (123 ppm) in its
C NMR spectrum. These peaks and assignments are
consistent with the existing literature (Gao et al. 2009). After gelation, the epoxide,
hydroxyl, and carbonyl peaks are virtually eliminated and an aliphatic carbon peak
(26 ppm) appears. The disappearance of the large peaks between 50 and 75 ppm in
the gel suggests that epoxide and hydroxyl groups are involved in the cross-linking
mechanism. Conversely, the emergence of the aliphatic carbon (sp
) peak suggests
that CH
- and/or CH
O- are present in the gel. The CH
- and CH
O- moieties
likely function as the cross-links that support the initial 3D GO network similar to
the cross-links formed in resorcinol-formaldehyde (RF) solgel chemistry to form
organic gels (Al-Muhtaseb and Ritter 2003; Pekala and Kong 1989).
spectra (Fig. 18b) for the sample after gelation also support the presence of CH
and CH
O- moieties with peaks at 0.9 and 3.1 ppm. The presence of CH
moieties is further supported by energy dispersive X-ray (EDX) analysis which
measured 11 at.% oxygen remaining in the initial aerogel. After pyrolysis, only the
carbon peak remains suggesting that the sp
carbon cross-links were thermally
converted to conductive sp
carbon junctions, again analogous to the carbonization
process that occurs during the pyrolysis of resorcinol-formaldehyde-based gels. The
H NMR spectrum (Fig. 18b) supports the conversion of the CH
- and CH
moieties with a virtual elimination of those peaks in the thermally treated sample.
Lastly, the reduction of carbon is conrmed by oxygen content of less than 2 at.%, as
determined by EDX, in the nal graphene assembly.
Fig. 18
C and
H NMR spectra for GO powder, GO after initial gelation, and graphene aerogel
(Reproduced from reference (Worsley et al. 2012) with the permission of the Royal Society of
Carbon Aerogels 25
The SEM images in Fig. 19 show that the 3D graphene monolith has a sheetlike
microstructure similar to that reported in other graphene assemblies. In particular, the
morphology resembles that of an RF-free graphene assembly reported to have a
surface area in excess of 1000 m
/g, presumably due to minimal thickness of
graphene sheets (Worsley et al. 2011b). Nitrogen porosimetry results for the material
are consistent with the morphology revealed by SEM. The nitrogen adsorption/
desorption isotherm shown in Fig. 20a is Type IV, indicative of a mesoporous
material. The observation of a Type 3 hysteresis loop (IUPAC classication) at
high relative pressure is consistent with other 3D graphene materials, but the
increased magnitude of the loop is indicative of a much larger pore volume than
those reported for other graphene assemblies. The BET surface area for this graphene
aerogel is 1314 m
/g or roughly half of the theoretical value expected for a single
graphene sheet. This extremely high surface area compared to assemblies made
using solgel chemistry (Worsley et al. 2010) suggests that layering/overlapping of
sheets has been signicantly reduced with the direct cross-linking approach. The
reduction in layering of sheets is also consistent with an XRD pattern that lacks a
strong (002) peak at ~28(graphite interlayer spacing). The pore size distribution
(Fig. 20b) shows that much of the pore volume (4.0 cm
/g) lies between 3 and
10 nm, with a peak pore diameter at 6 nm.
The mechanical behavior of the graphene aerogels revealed a mechanical behav-
ior qualitatively similar to that of a CNT aerogel. The graphene aerogel has a
Youngs modulus of 51 12 MPa, which is orders of magnitude higher than
those reported for other graphene aerogels (Tang et al. 2010; Xu et al. 2010;
Zhang et al. 2011; Sheng et al. 2011). In addition to being extraordinarily stiff, the
graphene aerogels exhibit super-compressive behavior with failure strains of
57 21 % and a complete recovery for lower strains. The failure stress is
10.4 3.9 MPa. These values of failure stress and strain are comparable to those
of CNT aerogels of the same density (100 mg/cc). These remarkable mechanical
properties can be attributed to the robustness and preponderance of sp
carbon cross-
Fig. 19 SEM images of the fracture surface of the graphene aerogel at (a) low and (b) high
magnication (Reproduced from reference (Worsley et al. 2012) with the permission of the Royal
Society of Chemistry)
26 M.A. Worsley and T.F. Baumann
links between graphene sheets, in addition to the excellent mechanical properties of
the graphene sheets themselves.
Bulk electrical conductivity of the graphene aerogel, evaluated by the four-probe
method, was measured at 100 S/m. This is consistent with carbon junctions cross-
linking graphene sheets (Worsley et al. 2010) and is orders of magnitude higher than
for graphene aerogels made via physical cross-links. Cyclic voltammetry (CV) was
used to characterize the energy storage capabilities of the graphene aerogels in
aqueous electrolyte (5 M KOH). At low scan rates, the CVs exhibit the typical
rectangular shape expected for pure double-layer capacitors like conventional carbon
aerogels, as well as CNT and graphene aerogels (Fig. 21a). Analysis of the CVs
measured at low scan rates reveals a maximum capacitance of 165 F/g. Remarkably,
the 500 μm thick graphene aerogel electrode was able to maintain greater than 50 %
of its maximum capacitance (89 F/g) up to 100 mV/s, indicating an exceptionally
fast charge/discharge capability. The graphene aerogel has a maximum energy
density of 27 Wh/kg and a maximum power density approaching 10 kW/kg
(Fig. 21b). Further optimization of the electrodes, such as using thinner electrodes
(e.g., 100 vs. 500 μm thickness), and electrolyte (e.g., inorganic vs. aqueous) could
push the power and energy densities to ~10
kW/kg and ~10
Wh/kg, respectively.
These observations illustrate the potential of graphene aerogels for energy storage
In summary, we have developed a graphene aerogel material that combines high
surface area, high electrical conductivity, and mechanical robustness. Our design
approach utilizes the functional groups native to GO as direct cross-linking sites.
These cross-links are then converted to sp
carbon as the GO is reduced to graphene
by thermal annealing, providing strong, conductive junctions between graphene
sheets. The resulting aerogel has large capacitance (165 F/g), high energy (27 Wh/
kg), and power density (10 kW/kg). In addition to energy storage applications, these
Fig. 20 (a) Nitrogen adsorption/desorption isotherm and (b) pore size distribution for the graphene
aerogel (Reproduced from reference (Worsley et al. 2012) with the permission of the Royal Society
of Chemistry)
Carbon Aerogels 27
3D graphene assemblies should nd application in other areas as well, including gas
storage, sensors, and catalysis.
Graphene Xerogels
These low-density nanoporous graphene structures exhibit electrical conductivities
and Youngs moduli as many as 10 orders of magnitude lower than those observed
for individual graphene sheets, which is not surprising given their high porosity. It is
a direct consequence of superlinear dependences of electrical and mechanical prop-
erties on the monolith density for porous materials (Pekala et al. 1990,1995;
Worsley et al. 2009a; Gross et al. 1992). A recent attempt to address this with higher
density assemblies have resulted in improvements in some mechanical properties,
but the electrical conductivity has remained relatively low and comparable to that of
low-density counterparts (Bi et al. 2012). Therefore, realizing macroscale 3D
graphene-based materials exhibiting the exceptional properties of graphene sheets
are a signicant challenge.
Here, a straightforward and relatively low-temperature method to realize macro-
scopic graphene xerogels with isotropic properties approaching those of graphene
sheets is described (Worsley et al. 2014). Essential elements to assembling a
macrostructure with such properties are (i) the development of strong covalent
links between graphene sheets that facilitate both electrical conductivity and struc-
tural reinforcement, (ii) sufcient restacking of graphene sheets to reach relatively
high densities of ~1 g/cm
, and (iii) an assembly that can be carbonized at relatively
low temperatures (~1000 C).
In order to form strong links between graphene sheets in graphene xerogels, the
present method is based on chemical cross-linking of individual sheets of graphene
oxide (GO) suspended in water. Such cross-linking involves various GO functional
groups (e.g., epoxide and hydroxide), yielding a reduced GO gel. Upon annealing at
Fig. 21 (a) CV scans and (b) Ragone plot for 500 mm thick graphene aerogel electrodes
(Reproduced from reference (Worsley et al. 2012) with the permission of the Royal Society of
28 M.A. Worsley and T.F. Baumann
about 1000 C the organic cross-links are reduced to sp
carbon cross-links. This
approach was recently demonstrated in the synthesis of exceptionally stiff and
electrically conductive low-density graphene aerogels (Worsley et al. 2012). The
restacking of graphene sheets in the graphene xerogel is achieved by simply drying
the reduced GO gel under ambient conditions instead of using a supercritical solvent
extraction method that preserves the low-density structure and minimizes sheet
restacking in graphene aerogels. The dried monoliths are then annealed to produce
carbon cross-links. Resultant graphene xerogels display isotropic bulk properties
such as electrical conductivities and Youngs moduli that are three to six orders of
magnitude higher than for any previously reported 3D graphene assembly as well as
exceeding the properties of isotropic graphite (
com/cgrawmaterials.html) that is nearly twice as dense.
Synthesis of graphene xerogels was carried out via gelation of a GO suspension
under basic conditions. Briey, in a typical synthesis, ultrasonication is used to
disperse 2 wt% GO in deionized water. Concentrated NH
OH is added (211 μl/g GO
suspension) to the suspension, which is then sealed and placed in an oven at 85 Cto
gel. After gelation, the reduced GO gel is washed rst in deionized water and then in
acetone. After that, the gel is allowed to dry under ambient conditions for at least
24 h, followed by annealing at 1050 C under nitrogen to yield the nal graphene
xerogel with a monolith density of ~1.0 g/cm
. This is an ~10-fold increase in the
density compared to graphene aerogels that have densities of ~0.10 g/cm
. Most
importantly, the randomly orientated self-assembly during gelation combined with
the capillary force-induced shrinkage leads to the formation of a completely isotropic
material. There are few limits on the shapes and sizes of the graphene xerogels
produced since, in addition to being readily mechanically machined, the graphene
xerogel conforms to the mold in which the initial reaction mixture is placed (Fig. 22).
A large difference in the microstructure of graphene xerogels and graphene
aerogels is illustrated by SEM images in Fig. 22. Graphene aerogels are made of
randomly interconnected graphene sheets with a minimal restacking (Fig. 1b,c)
resulting in ne sheetlike features and a large porosity of ~95 % (Worsley
et al. 2010,2012; Wang and Ellsworth 2009; Xu et al. 2010; Bai et al. 2011; Chen
and Yan 2011; Sui et al. 2011; Yin et al. 2012). In contrast, no ne sheetlike features
are distinguishable in the graphene xerogel in Fig. 22d,e. The SEM shows that the
graphene xerogel is homogenous and contains much less porosity than the graphene
aerogel. This lower porosity contributes to the materials density being roughly half
of the density for single crystalline graphite made of perfectly stacked graphene
sheets (2.2 g/cm
) compared to the ultralow-density graphene aerogel. This apparent
porosity of the graphene xerogel is consistent with the measured Brunauer-Emmett-
Teller (BET) surface area of 69 m
/g determined by nitrogen porosimetry.
TEM gives further details concerning the sheet assembly in the graphene xerogel
and illustrates one of the key differences between the graphene xerogel and graphite:
the use of chemical cross-linking versus Van der Waals (VdW) forces to govern
sheet-to-sheet interactions. The TEM images of commercial graphite show large
domains of highly ordered, perfectly stacked graphene sheets (Fig. 22f). This perfect
Carbon Aerogels 29
stacking is only possible if the sheets are loosely associated with each other (VdW)
and are able to slide until this nal conguration is achieved. In contrast, no regimes of
perfectly stacked graphene sheets are observed in the graphene xerogel (Fig. 22g,h).
The sheets are closely associated with each other but possess a spectrum of curvature
and orientations. We suggest that this imperfect or partial stacking is due to the fact
that, like the graphene aerogel, the graphene sheets are cross-linked at various points
by sp
carbon bonds (Worsley et al. 2011b,2012). These strong chemical cross-links
make it impossible for the sheets to slide into the perfectly stacked layers observed in
graphite. The incomplete stacking also contributes to the lower density of the graphene
xerogel compared to commercial graphite.
Despite having a density of close to half of that of perfectly stacked graphene
sheets (2.2 g/cm
) or even commercial grade graphite (1.61.9 g/cm
), the mechan-
ical properties of the graphene xerogel (randomly stacked, cross-linked sheets)
exceed those of graphite. Nanoindentation results show that the graphene xerogel
exhibits a Youngs modulus of ~10 GPa and a failure stress of ~1.2 GPa independent
of the loading direction. These exceed respective nanoindentation values of 8.5 and
Fig. 22 (a) Photograph of a reduced GO wet gel before drying (left), after supercritical drying
(middle), and after ambient drying (right). For reference, the wet gel is in a 20 ml vial. Inset shows a
xerogel cylinder (cast), prism (machined), and pyramid (machined). SEM images of fracture
surfaces of the aerogel (b,c) and xerogel (d,e) at low (b,d) and high (c,e) magnication. (f)
TEM images of commercial graphite at high magnication. Inset is a zoom-in of white box area and
is 10 nm in width. (g) TEM image of xerogel at high magnication. Inset is a zoom-in of white box
area and is 10 nm in width. (h) Low magnication TEM image of the xerogel (Reprinted by
permission of John Wiley & Sons, Inc)
30 M.A. Worsley and T.F. Baumann
0.4 GPa (also independent of loading direction) for commercial grade graphite
( with a much larger density of
1.7 g/cm
. Uniaxial compression results also show that the graphene xerogel has a
high compressive strength of ~200 MPa compared to only ~98 MPa for the com-
mercial graphite material. The measured compressive strength for commercial
graphite is consistent with reported values of 40193 MPa, depending on the
material density (1.61.9 g/cm
) and defect content (
graphite-rods/=kkbdgu;;http://www. The much-improved mechanical properties
of the graphene xerogel can be attributed to its unique nanostructure. As seen in
the TEM images, though there is signicant restacking of the graphene sheets, there
is no preferred orientation, resulting in isotropic mechanical response. Furthermore,
instead of relying on weak VdW forces between sheets as in commercial graphite,
the graphene xerogel graphene sheets are cross-linked with strong sp
bonds. These
chemical cross-links allow for much stronger structural reinforcement between
graphene sheets than found in traditional graphite.
Considering that Youngs modulus depends superlinearly on monolith density, ρ,
for nanoporous materials with an exponent, n, in the range of 24, it was that the
modulus of the graphene xerogel scales with density as expected for CNT- or
graphene-based materials with n=2.5 (Worsley et al. 2014). This suggests that if
the density of the graphene xerogel could be further increased to ~2 g/cm
, the
modulus would come within an order of magnitude of that for an individual
atomically perfect graphene sheet (10
MPa). The scaling of the modulus to a
value lower than that of an atomically perfect graphene sheet probably stems from
random orientation and curvature of graphene sheets as well as lattice defects.
Nevertheless, it is quite remarkable that the graphene xerogel exhibits properties
so close to those expected based on the density scaling of an assembly of atomically
perfect graphene sheets.
The bulk electrical conductivity of the graphene xerogel, determined via the four-
probe method, is 1750 S/cm, independent of probe orientation, which, to our
knowledge, is larger than that observed for any graphene aerogel reported by several
orders of magnitude (Worsley et al. 2012; Bi et al. 2012). Even high-density
graphene structures that have been annealed under similar conditions exhibit far
lower conductivities than the graphene xerogel (Bi et al. 2012). Similar to the scaling
of Youngs modulus discussed above, electrical conductivity, σ, scales superlinearly
with the material density (σ~ρ
). It was observed that, despite pronounced
restacking of graphene sheets revealed by an increased density and combined
SEM, TEM, and BET observations, the graphene xerogel follows the scaling law
for CNT- and graphene-based materials, extrapolating to a conductivity of ~10
cm at ~2 g/cm
(Worsley et al. 2014). This result is approximately two orders of
magnitude less than what is observed for a perfect graphene sheet, highlighting how
sensitive the electrical conductivity is to defects, sheet curvature, sheet stacking, and
orientation (in-plane vs. through-plane). A scaled graphene xerogel conductivity of
S/cm is, however, consistent with experimental conductivity values for indi-
vidual graphene sheets thermally reduced from GO (Wu et al. 2009) and is an order
Carbon Aerogels 31
of magnitude higher than that of commercial graphite (
graphite-rods/=kkbdgu;;http://www. This improvement in electrical conductivity,
as was the case with the mechanical properties, can be attributed to the unique
nanostructure of the graphene xerogel. Conductive sp
bonding between graphene
sheets in all directions lowers the sheet-to-sheet resistance compared to traditional
graphite, which is limited to higher resistance VdW sheet-to-sheet bonding.
In summary, we have presented a straightforward method to realize high-density
graphene-based xerogels with isotropic mechanical and electronic transport proper-
ties approaching those of individual graphene sheets. The synthesis strategy
involved direct cross-linking of graphene sheets via the functional groups in
graphene oxide and sufcient restacking of the graphene sheets to increase the
material density. The nature of the synthesis requires a much lower temperature
than for the synthesis of commercial isotropic graphite and facilitates fabrication
with few limits on the size or shape. We anticipate that the development of 3D bulk
macrostructures retaining properties of individual graphene sheets will both expand
and accelerate the commercialization of graphene-based technologies and products.
Various solgel methods for realizing carbon-based aerogels have been reviewed. In
addition to traditional RF-derived carbon aerogels, the methods for synthesizing
graphene and CNT-based aerogels consist of roughly two categories. The rst
involves using RF solgel chemistry to induce gelation of a suspension of the desired
carbon nanomaterials (e.g., CNT, graphene, GO). The aqueous suspension is pre-
pared by ultrasonication and is used with small concentrations of RF such that the
RF-derived organic nanoparticles only coat and cross-link the carbon nanomaterials.
After thermal reduction, the RF-derived carbon both mechanically reinforces and
provides conductive bridges to the carbon nanomaterial network.
The second method involves the use of the chemical functionality present on the
carbon nanomaterial (e.g., GO) to induce gelation of the suspension. Here, gelation
of the aqueous suspension is triggered by a change in pH resulting in a reaction
between carbon nanomaterials themselves to form direct cross-links. And after a
thermal reduction, these direct cross-links are converted to sp
carbon bridges
between the carbon nanomaterials network. These are the methods most commonly
applied to solgel synthesis of modern CNT and graphene aerogels, which have
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36 M.A. Worsley and T.F. Baumann
... The traditional methods to create amorphous carbon aerogels include sol-gel processes, freeze drying of carbon suspensions. But today, the typical carbon aerogel consists of a fine 3D network of CNT with an increased Young's modulus i.e. super compression and elasticity in addition to its existing mechanical properties 8 . It has enhanced its mechanical and transport properties, despite its delicate appearance, supported entirely by Van der Waals forces between adjacent CNT. ...
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Everyone wants the lightest materials to their use. Lighter materials are the most important parameter for a dynamic system. This chapter describes all the types of latest ultralight materials having a density of <10 mg/cm³. Aerogel, aerographite, aerographene, 3D graphene, carbyne, microlattice, and foam come under ultralight material. Classification, fabrication, properties, and applications of each of the ultralight material has been described here extensively.
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A sensitive sandwich-model immunoassay designed for simultaneous electrochemical measurement of alphafetoprotein and carcinoembryonic antigen handling multiple-label strategy. The sandwich-model immunosensor was constructed by assembling ethylenediamine-MWCNT aerogels (EDA-CAGs), AuNPs and capture antibodies (Cp-Ab1) on the screen-printed carbon electrode (SPCE). The prepared EDA-CAG textures were analyzed using Fourier transform infrared (FT-IR), Scanning Electron Microscopy (SEM) and Energy Dispersive X-Ray Spectroscopy (EDX). In this protocol, AuNPs decorated EDA-CAG-carried Thionine (Thi) and AuNPs decorated EDA-CAG-carried Safranine O (SfO) (denoted as AuNP-Thi-EDA-CAG and AuNP-SfOEDA- CAG, respectively) were properly exploited as distinguishable signal labels, which were used to attach detection antibodies (Ab2) (anti-AFP2 and anti-CEA2), respectively. When two tumor antigens were present, a square wave voltammetry (SWV) scan displayed two well-separated signals, each signal indicated one target antigen. Observed results introduced that the square wave voltammetric peak current exhibited a good linear relationship to logarithm concentration in the ranges from 0.005 to 1.0 ng/mL for both analytes. The LoDs were 0.0015 and 0.0010 ng/mL (at signal/noise S/N = 3) for AFP and CEA, respectively. The designed sandwich- type immunosensor was applied to real serum sample analysis.
This chapter basically summarizes progress in the design and fabrication of three-dimensional graphene-based nanocomposites for promising solicitations. The graphene nanofoams (or nanosponges) have been prepared using several advance approaches. These nanostructures have shown high specific surface area, ultrahigh porosity, flexibility, and outstanding electrical conductivity, thermal transportation, and mechanical robustness. The three-dimensional macroscopic graphene nanostructures have been reinforced in polymer matrices to attain superior structural and physical properties of the resulting nanocomposite. Polymer/graphene nanofibers have also been studied in the nanofoam architectures. The practical applications of hierarchical graphene nanofoams were observed in the fields of batteries, supercapacitors, fuel cells, solar cells, radiation shielding, and electronic devices.
The increasing environmental pollution demands sustainable remediation materials for cleaning both air and water. In the past decade, carbon aerogels have drawn considerable interest due to their pertinent properties for environmental remediation. Properties like chemical stability, low density, high porosity, abundant pore structure, large pore volume, high specific surface area, and adaptive surface chemistry make them suitable candidates for CO2 capture, oil-water separation, volatile organic compounds (VOCs) adsorption, organic dyes pollutants, and heavy metal ion adsorption. This book chapter details the various preparation methods of carbon aerogels from activated carbon, graphene, carbon nanotube, and composite aerogel along with their applications in CO2 capture, VOC removal, oil-water separation, and heavy metal ion removal.
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Because of pharmaceutical-emerging contaminants in water resources, there has been a significant increase in the antibiotic resistance in bacteria. Therefore, the removal of antibiotics from water resources is essential. Various antibiotics have been greatly studied using many different carbon-based materials including graphene-based hydrogels and aerogels. In this study, carbon aerogels (CAs) were synthesized from waste paper sources and their adsorption behaviors toward three antibiotics (hygromycin B, gentamicin, and vancomycin) were investigated, for which there exist a limited number of reports in the literature. The prepared CAs were characterized with scanning electron microscopy, transmission electron microscopy, X-ray photoelectron spectroscopy, and micro-computerized tomography (μ-CT). According to the μ-CT results, total porosity and open porosity were calculated as 90.80 and 90.76%, respectively. The surface area and surface-to-volume ratio were found as 795.15 mm² and 16.79 mm–1, respectively. The specific surface area of the CAs was found as 104.2 m²/g. A detailed adsorption study was carried out based on different pH values, times, and analyte concentrations. The adsorption capacities were found as 104.16, 81.30, and 107.52 mg/g for Hyg B, Gen, and Van, respectively. For all three antibiotics, the adsorption behavior fits the Langmuir model. The kinetic studies showed that the system fits the pseudo-second-order kinetic model. The production of CAs, within the scope of this study, is safe, facile, and cost-efficient, which makes these green adsorbents a good candidate for the removal of antibiotics from water resources. This study represents the first antibiotic adsorption study based on CAs obtained from waste paper.
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In this study, synthesis and characterization of thick carbon nanotube (CNT) electrodes with impressive electrical conductivity, mechanical properties, and large surface area is presented. The electrodes are made by effectively cross-linking the CNTs with graphitic carbon particles. The random network of cross-linked CNTs results in a bulk conductivity of >60 S/cm, large Young's modulus (similar to 1.2 GPa), and high surface area (>500 m(2)/g). Electrochemical characterization of these thick CNT electrodes (100-1000 mu m) shows a capacitance of 90 F/g (40 F/cm(3)) in an aqueous electrolyte (5 M KOH). Furthermore, much of this capacitance is maintained at fast scan rates up to 1000 mV/s. This high rate performance leads to a simultaneous display of large energy densities (> 10 Wh/kg) and power densities (>40 kW/kg) for the thick, binderless CNT electrode.
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The remarkable electrical and thermal conductivities of isolated carbon nanotubes have spurred worldwide interest in using nanotubes to enhance polymer properties. Electrical conductivity in nanotube/polymer composites is well described by percolation, where the presence of an interconnected nanotube network corresponds to a dramatic increase in electrical conductivity ranging from 10−5 S/cm to 1 S/cm. Given the high aspect ratios and small diameters of carbon nanotubes, percolation thresholds are often reported below 1 wt% although nanotube dispersion and alignment strongly influence this value. Increases in thermal conductivity are modest (∼3 times) because the inter facial thermal re sis tance between nanotubes is considerable and the thermal conductivity of nanotubes is only 104 greater than the polymer, which forces the matrix to contribute more toward the composite thermal conductivity, as compared to the contrast in electrical conductivity, >1014. The nanotube network is also valuable for improving flame-retardant efficiency by producing a protective nanotube residue. In this ar ticle, we highlight published research results that elucidate fundamental structure–property relationships pertaining to electrical, thermal, and/or flammability properties in numerous nanotube-containing polymer composites, so that specific applications can be targeted for future commercial success.
The remarkable electronic and structural properties of semiconductor surfaces and interfaces result from the existence of surface and interface states, respectively. Surface states on clean surfaces originate from dangling bonds and on adsorbate-covered surfaces from bonds between adsorbate and semiconductor-surface atoms. At abrupt metal-semiconductor interfaces, the wavefunctions of those metal electrons, which energetically overlap the semiconductor band gap, decay exponentially into the semiconductor. These tails represent metal-induced interface states. This concept also applies to semiconductor het-erostructures and semiconductor-insulator interfaces. Surface and interface states above the bulk valence-band maximum may become charged. Surface charge neutrality then requires the existence of space-charge layers which penetrate from the surface or interface into the semiconductor.
The preparation of graphitic oxide by methods described in the literature is time consuming and hazardous. A rapid, relatively safe method has been developed for preparing graphitic oxide from graphite in what is essentially an anhydrous mixture of sulfuric acid, sodium nitrate and potassium permanganate.
We have demonstrated for the first time the creation of graphene oxide and graphene-based carbon aerogels by freeze-drying of graphite oxide/water dispersion. The graphene oxide aerogel consists of sheet-like nanocarbon structures, and could be significantly reinforced by small amounts of water soluble polymers. Controlled thermal treatments reduce graphene oxide to graphene and restore the electrical conductivity of the graphene aerogel. The graphene aerogels are electrically conducting, with extremely high surface-area-to-volume ratio, could outperform conventional carbon aerogels for applications such as sensors, actuators, high performance polymer nanocomposites, and electrochemical applications such as porous electrodes for batteries, fuel cells, and supercapacitors.
Compressive graphene aerogels were obtained by the one-step reduction and self-assembly of graphene oxide with ethylenediamine and then freeze-drying. The aerogels hold good compressibility, variable electrical resistance and fire-resistance. The high porosity with a hydrophobic nature, allows the aerogels to absorb different organic liquids, and the absorption–squeezing process has been demonstrated for oil collection.
Rabenschwarzes Leichtgewicht: Dreidimensionale Makrostrukturen (siehe Bild) entstehen aus Graphenoxid‐Schichten und Edelmetallnanokristallen (Au, Ag, Pd, Ir, Rh, Pt usw.). Ungeachtet seiner überaus geringen Dichte (ca. 0.03 g cm−3) verfügen Zylinder wie der abgebildete über hervorragende mechanische Eigenschaften. Als Festbettkatalysator vermittelten sie eine Heck‐Reaktion mit annähernd vollständiger Selektivität und Umwandlung.
Realization of macroscale three-dimensional isotropic carbons that retain the exceptional electrical and mechanical properties of graphene sheets remains a challenge. Here, a method for fabricating graphene-derived carbons (GDCs) with isotropic properties approaching those of individual graphene sheets is reported. This synthesis scheme relies on direct cross-linking of graphene sheets via the functional groups in graphene oxide to maximize electronic transport and mechanical reinforcement between sheets and the partial restacking of the sheets to increase the material density to about 1 g cm-3. These GDCs exhibit properties 3–6 orders of magnitude higher than previously reported 3D graphene assemblies.