Embed Size (px)
Citation preview
ADVANCED FABRICATION AND MULTIFUNCTIONAL
PROPERTIES OF MORPHOLOGY-CONTROLLED
GRAPHENE AEROGELS AND THEIR COMPOSITES
FAN ZENG
(B. Eng, Harbin Institute of Technology)
A THESIS SUBMITTED
FOR THE DEGREE OF DOCTOR OF PHILOSOPHY
DEPARTMENT OF MECHANICAL ENGINEEING
NATIONAL UNIVERSITY OF SINGAPORE
2015
i
ACKNOWLEDGEMENTS
I would like to convey my deepest and sincerest appreciation to my
supervisors, Asst. Prof. Duong Hai Minh and Assoc. Prof. Christina Lim Y. H.,
for their invaluable guidance and constant support during my whole PhD
candidature. The energy and enthusiasm that they have for research never fail to
inspire me.
I extend my sincere thanks to Assoc. Prof. Brian L. Wardle and all group
members at Massachusetts Institute of Technology (MIT) for their warm
hospitality and valuable suggestions for my research during my visit to
NECSTlab.
I am extremely grateful to my fellow group members, Dr. Nguyen Truong
Son, Mr. Gong Feng, Ms. Feng Jingduo, Mr. Cheng Hanlin, Ms. Liu Peng, Mr.
Tran Quyet Tran, Mr. Daniel Tng Z. Y., Ms. Clarisse Lim X. T., and Mr. Daryl
Lim, for their devoted help and cooperation in my research work. I am also
thankful to Asst. Prof. Chua Kian Jon, Dr. Zhao Xing, Mr. Sun Bo, and Asst.
Prof. Amy Marconnet at Purdue University for their support regarding thermal
conductivity measurements.
I also thank the technical staff in Materials Laboratory, Mr. Thomas Tan,
Mr. Ng Hongwei, Mr. Abdul Khalim Bin Abdul, and Mr. Juraimi B. Madon,
and the technical staff in the Energy Conversion Laboratory, Mr. Tan Tiong
Thiam, for their superior and professional technical support.
ii
Finally, I gratefully acknowledge the Start-up Grant R-265-000-361-133,
SERC 2011 Public Sector Research Funding (PSF) Grant R-265-000-424-305,
and the China Scholarship Council (CSC) for providing financial support.
iii
TABLE OF CONTENTS
ACKNOWLEDGEMENTS ............................................................................. i
TABLE OF CONTENTS .............................................................................. iii
SUMMARY .................................................................................................... vii
LIST OF TABLES .......................................................................................... ix
LIST OF FIGURES ......................................................................................... x
CHAPTER 1: Introduction .......................................................................... 1
1.1 Aerogels .............................................................................................. 1
1.1.1 Background ............................................................................ 1
1.1.2 Carbon-based aerogels ........................................................... 2
1.1.3 Drying techniques of aerogels ............................................... 4
1.1.4 General properties and applications of aerogels .................... 6
1.2 Carbon-based nanocomposites .............................................................. 8
1.3 Objectives of thesis ............................................................................... 9
1.4 Organization of thesis ......................................................................... 11
CHAPTER 2: Literature Review .............................................................. 13
2.1 Graphene ............................................................................................. 13
2.1.1 Introduction .......................................................................... 13
2.1.2 Structure of graphene ........................................................... 14
2.1.3 Synthesis methods of graphene ............................................ 15
2.1.4 Properties of graphene ......................................................... 24
2.1.5 Potential applications of graphene ....................................... 28
2.2 Graphene aerogels (GAs) .................................................................... 29
iv
2.2.1 Introduction .......................................................................... 29
2.2.2 Graphene oxide (GO) ........................................................... 31
2.2.3 Structure of GAs .................................................................. 34
2.2.4 Synthesis methods of GAs ................................................... 35
2.2.5 Properties of GAs ................................................................. 39
2.2.6 Potential applications of GAs .............................................. 42
2.3 Graphene–carbon nanotube (CNT) hybrid aerogels ........................... 44
2.3.1 Introduction .......................................................................... 44
2.3.2 Structure of graphene–CNT hybrid aerogels ....................... 44
2.3.3 Synthesis methods of graphene–CNT hybrid aerogels ........ 45
2.3.4 Properties and potential applications of graphene–CNT
hybrid aerogels ..................................................................... 46
2.4 Graphene–polymer nanocomposites ................................................... 47
2.4.1 Introduction .......................................................................... 47
2.4.2 Structure of graphene–polymer nanocomposites ................. 48
2.4.3 Synthesis methods of graphene–polymer nanocomposites .. 49
2.4.4 Properties of graphene–polymer nanocomposites ............... 50
2.4.5 Potential applications of graphene–polymer nanocomposites
.............................................................................................. 54
CHAPTER 3: Experimental Section ......................................................... 56
3.1 Materials ............................................................................................. 56
3.2 Experimental techniques ..................................................................... 57
3.2.1 Synthesis of graphene oxide (GO) ....................................... 57
3.2.2 Synthesis of graphene aerogels (GAs) ................................. 58
v
3.2.3 Synthesis of graphene–CNT hybrid aerogels ...................... 59
3.2.4 Synthesis of GA–poly (methyl methacrylate) (PMMA)
nanocomposites .................................................................... 59
3.3 Characterization .................................................................................. 60
3.3.1 X-ray diffraction (XRD) ...................................................... 60
3.3.2 Scanning electron microscopy (SEM) ................................. 60
3.3.3 Physical adsorption/desorption of nitrogen (BET) .............. 61
3.3.4 Electrical conductivity measurement ................................... 63
3.3.5 Thermal conductivity measurement ..................................... 64
3.3.6 Thermogravimetric analysis (TGA) ..................................... 67
3.3.7 Vickers microhardness measurement .................................. 67
CHAPTER 4: Morphology Control of Graphene Aerogels .................... 69
4.1 Introduction ......................................................................................... 69
4.2 Nanostructured control of graphene aerogels (GAs) .......................... 70
4.2.1 Effects of chemical compositions and synthesis conditions 72
4.2.2 Effects of reducing agents .................................................... 76
4.2.3 Effects of CNTs ................................................................... 77
4.3 Conclusions ......................................................................................... 79
CHAPTER 5: Multifunctional Properties of Graphene Aerogels .......... 80
5.1 Introduction ......................................................................................... 80
5.2 Electrical property of graphene aerogels (GAs) ................................. 81
5.2.1 Effects of chemical compositions and synthesis conditions 81
5.2.2 Effects of reducing agents .................................................... 83
5.2.3 Effects of CNTs ................................................................... 85
vi
5.3 Thermal stability of graphene aerogels (GAs) .................................... 86
5.3.1 Effects of chemical compositions and synthesis conditions 86
5.3.2 Effects of reducing agent and CNTs .................................... 87
5.4 Thermal conductivity of graphene aerogels (GAs) ............................. 88
5.4.1 Effects of chemical compositions ........................................ 88
5.4.2 Effects of thermal annealing ................................................ 93
5.4.3 Prediction of thermal boundary resistance and thermal
conductivity of graphene nanosheets within GAs ............... 94
5.5 Conclusions ......................................................................................... 96
CHAPTER 6: Advanced Multifunctional Graphene Aerogel–Poly
(methyl methacrylate) Composites: Experiments and Modelling ............. 98
6.1 Introduction ......................................................................................... 98
6.2 Structure of GA–PMMA nanocomposites .......................................... 99
6.3 Electrical property of GA–PMMA nanocomposites ......................... 101
6.4 Thermal stability of GA–PMMA nanocomposites ........................... 103
6.5 Thermal property of GA–PMMA nanocomposites .......................... 104
6.6 Microhardness of GA–PMMA nanocomposites ............................... 106
6.7 Conclusions ....................................................................................... 109
CHAPTER 7: Conclusions and Recommendations ............................... 111
7.1 Conclusions ........................................................................................ 111
7.2 Recommendations .............................................................................. 114
REFERENCES ............................................................................................. 116
LIST OF PUBLICATIONS ........................................................................ 160
vii
SUMMARY
Graphene aerogels (GAs) having large surface area property of aerogels
and excellent multifunctional properties of graphene nanosheets, can be
promising candidates for energy storage devices and light-weight
nanocomposites. For this research, a cost-effective method has been developed
to synthesize self-assembled GAs from graphite at a low temperature range
(45–95 oC) through modified Hummers’ and chemical reduction methods
using various reducing agents (L-ascorbic acid, NaHSO3 and HI). The effects
of synthesis conditions, reducing agents, thermal annealing processes and
carbon nanotubes (CNTs) on the morphology, electrical and thermal properties
of the as-prepared GAs are quantified comprehensively. The GA
nanostructures are well controlled and have a large surface area of up to 577
m2/g. The CNT inclusions and annealing processes can enhance the electrical
conductivity of the as-prepared GAs by up to five times, as measured via a
two-probe method. For the first time, a comparative infrared microscopy
technique has been successfully developed to measure the thermal conductivity
of GAs, and the GAs having 0.67–2.50 vol. % graphene are measured to be
0.12–0.36 W/m·K accordingly.
For light-weight but strong material development, GA–poly (methyl
methacrylate) (PMMA) nanocomposites are further developed by backfilling
PMMA into the pores of the GAs. Due to uniform distribution of the graphene
nanosheets in the PMMA matrix, the GA–PMMA nanocomposites exhibit
significant enhancements of electrical conductivity (0.16–0.86 S/m),
viii
microhardness (303.6–462.5 MPa), and thermal conductivity (0.35–0.70
W/m·K) over pure PMMA and the graphene–PMMA nanocomposites prepared
by traditional powdery dispersion methods. The developed GAs and
GA–PMMA nanocomposites in this thesis can be applied in aerospace,
composite electrodes, biosensors, and energy harvest and storage devices.
ix
LIST OF TABLES
Table 1-1 Identification of aerogel properties and features, with their
applications [18]
. (Reproduced with permission from Ref. 18. Copyright (1998)
Elsevier.) ............................................................................................................ 7
Table 2-1 A comparison of strengths and drawbacks for the methods commonly
applied to synthesize graphene [68]
. (Reproduced with permission from Ref. 68.
Copyright (2010) Elsevier.) ............................................................................. 15
Table 4-1 Nanostructure control of the GAs. .................................................. 73
Table 4-2 Effects of reducing agent and addition of CNTs on morphology of
the GAs. ........................................................................................................... 76
Table 5-1 Morphological effects on electrical property of the GAs. .............. 82
Table 5-2 Effects of reducing agent and addition of CNTs on electrical
conductivity and thermal stability of the GAs. ................................................ 83
Table 5-3 Synthesis conditions, density, volume fraction and thermal
conductivity of the GAs. .................................................................................. 90
Table 6-1 Synthesis conditions, graphene volume fraction, and multifunctional
properties of GA–PMMA nanocomposites. .................................................. 100
x
LIST OF FIGURES
Figure 1.1 Application examples of carbon-based aerogels. (a) Representation
of an electric double-layer capacitor (EDLC); (b) Three-dimensional N-doped
graphene aerogel-supported Fe3O4 nanoparticles as efficient cathode catalysts
for the redox reaction [20]
. (Reproduced with permission from Ref. 20,
Copyright (2012) American Chemical Society.) ............................................... 3
Figure 1.2 Principles of three drying methods (evaporation, freeze drying, and
CPD) to remove liquid from hydrogels in a generic phase diagram. (a) Direct
evaporation from liquid phase to gas phase causes collapse to the gel network
due to the effect of surface tension. (b) Freeze drying is a common method to
create aerogels. During this process, the samples are fast frozen and then placed
in a vacuum under raised temperature to allow the ice to sublimate. (c) In the
CPD process, the temperature and pressure of liquid carbon dioxide are raised
above its critical point to form a supercritical fluid. By slowly releasing the
pressure, the hydrogels are dried to aerogels with well-preserved networks. The
CPD process is selected as the drying method in this thesis work. Adapted from
the thesis work of Mateusz B. Bryning [33]
. ....................................................... 6
Figure 2.1 Graphene is the building block of all carbon materials, e. g. 0D
fullerenes, 1D CNTs, and 3D graphite [24]
. (Reproduced with permission from
Ref. 24. Copyright (2007) Macmillan Publisher Ltd.) ..................................... 13
Figure 2.2 Scheme of the crystal structure, Brillouin zone, and dispersion
spectrum of graphene [52]
. (Reproduced with permission from Ref. 52.
Copyright (2010) Wiley–VCH.) ...................................................................... 15
Figure 2.3 Mechanically exfoliated (a) multi-layer, (b) few-layer and
monolayer graphene. Yellow colors indicate multi-layer graphene of thickness
over 100 nm, while blue, darker, and lighter purple indicate few-layer and
monolayer graphene [68]
. (Reproduced with permission from Ref. 68. Copyright
(2010) Elsevier.) .............................................................................................. 18
Figure 2.4 (a) A SEM image of a monolayer graphene suspended on an array of
circular holes, scale bar: 3μm; (b) A scheme of nanoindentation on the
membrane of graphene [54]
. (Reproduced with permission from Ref. 54.
Copyright (2008) AAAS.) ............................................................................... 26
Figure 2.5 Two sub-models of the Lerf-Klinowski model with (top) and
without (bottom) the presence of carboxylic acid groups on the periphery of the
graphitic basal plane of GO [162, 163, 169]
. (Top: Reproduced with permission from
xi
Ref. 162. Copyright (1998) American Chemical Society. Bottom: Reproduced
with permission from Ref. 163. Copyright (1998) Elsevier.) .......................... 32
Figure 2.6 Schematic of hydrogen bonding formed between water (H2O) and
the oxygen-containing functional groups of GO [169]
. (Reproduced with
permission from Ref. 169. Copyright (2010) The Royal Society of Chemistry.)
.......................................................................................................................... 33
Figure 2.7 (a) A typical SEM image which shows the structure of GAs. (b)
AFM images of typical building blocks of the GAs [157]
. (Reproduced with
permission from Ref. 157. Copyright (2011) Elsevier.) .................................. 35
Figure 2.8 SEM images showing the network structures of graphene–CNT
hybrid aerogels [157]
. (Reproduced with permission from Ref. 157. Copyright
(2011) Elsevier.) .............................................................................................. 45
Figure 2.9 Relationship between d spacing of graphene nanosheet (GNS)
families and graphite [210]
. Addition of CNTs and fullerene (C60) was found to
effectively increase the interlayer spacing of graphene. (Reproduced with
permission from Ref. 210. Copyright (2008) American Chemical Society.) .. 46
Figure 3.1 A Solartron 1260+1287 electrochemical system for electrical
property testing. ............................................................................................... 63
Figure 3.2 (a) Set-up schematic of thermal conductivity measurement using the
comparative infrared thermography technique. (b) Temperature distribution and
linear best fit curves for the 3-layered stack, which consists of a GA sample
sandwiched between two amorphous quartz layers, scale: 133 μm/ pixel. (c)
Heat flux as a function of temperature gradient in the GA region at several
power levels. During the measurement, the heat flux is typically below
11.25kW/m2. .................................................................................................... 66
Figure 4.1 (a) Aqueous suspension of GO and a graphene hydrogel and (b) A
graphene aerogel. ............................................................................................. 71
Figure 4.2 XRD patterns of graphite, graphite oxide and graphene aerogel. . 72
Figure 4.3 SEM images of graphene aerogels synthesized under different
conditions: (a) 2 mg GO/ml, 70 °C, 40 h; (b) 12 mg GO/ml, 70 °C, 40 h; (c) 2
mg GO/ml, 45 °C, 40 h; (d) 2 mg GO/ml, 95 °C, 40 h; (e) 2 mg GO/ml, 70 °C, 5
h. Compared with image (a), more graphene nanosheets can be observed in
image (b). The GA sample in image (c) shows a more loose structure than those
xii
in image (a) and image (d). Image (e) also shows a less dense network than that
in image (a). ..................................................................................................... 73
Figure 4.4 A typical N2 adsorption/desorption isotherm of a GA. ................. 74
Figure 4.5 (a) A digital photo of a graphene–CNT hybrid aerogel; (b) and (c)
SEM images of a graphene–CNT hybrid aerogel at different magnifications. 78
Figure 5.1 I–V plot of the GA obtained from 12 mg/ml GO suspension at 70 °C
for 40 h. ............................................................................................................ 81
Figure 5.2 Effects of initial GO concentrations on electrical conductivity and
surface area of GAs. ......................................................................................... 83
Figure 5.3 Effect of reducing agents and thermal annealing on (a) electrical
conductivity and (b) thermal stability of the GAs. .......................................... 84
Figure 5.4 XRD patterns of the GAs reduced by different reducing agents with
and without thermal annealing. ........................................................................ 85
Figure 5.5 Electrical conductivity and DTA peak temperature for
graphene–CNT hybrid aerogels with various CNT concentrations. ................ 86
Figure 5.6 Thermal stability of GAs at (a) 45 °C and (b) 95 °C. .................... 87
Figure 5.7 SEM images of the GAs synthesized under different conditions:
(a)–(h) represent GA1–GA8 respectively. All images present interconnected
3D networkstructures. With annealing, images (e)–(h) show denser networks
than those in images (a)–(d) at the same GO concentration. ........................... 89
Figure 5.8 Effects of GO concentrations and thermal annealing on thermal
conductivities of the GAs................................................................................. 91
Figure 5.9 XRD patterns of the GA4–5, 7 and 8. Broad peaks indicate the
non-crystal aerogel stucture. Compared with GA7, the peak of GA8 shifts right
and indicates the smaller interspacing between graphene nanosheets. ............ 94
Figure 5.10 The thermal condictivity of the GAs as a function of graphene
volume fraction. Best fits from rule of mixtures (dashed line) and EMT (orange
dots) are shown. For the rule of mixtures, the slope of the least-squares best fits
to the experimental data is the (kG-kair). For EMT, the values calculated assume
a kG=30.2 W/m·K, a value which is obtained by manually fitting the data. .... 95
xiii
Figure 6.1 Digital images of (a) GA and (b) GA–PMMA nanocomposite and
FE-SEM images of (c) the porous GA without PMMA infiltrated and (d) the
cross section of the GA–PMMA nanocomposite........................................... 100
Figure 6.2 (a) The electrical conductivity (black dots) of GA–PMMA
nanocomposites as a function of graphene volume fraction at room temperature.
The dashed blue line is the best fit from the power law relation 𝜎𝑐 =𝜎𝑓[(𝜑 − 𝜑𝑐)/(1 − 𝜑𝑐)]𝑡 with a nonlinear least-squares algorithm, where the
fitted values are σf=2543.8, φc=0 and t=2.18, respectively. (b) Comparative
study of electrical conductivity between this work and previously reported
graphene–PMMA nanocomposites synthesized via in situ polymerization. For
comparison purposes, the values of graphene loading reported in wt. % have
been converted to vol. %. ............................................................................... 102
Figure 6.3 Thermogravimetric curves of PMMA and GA–PMMA
nanocomposites. ............................................................................................. 104
Figure 6.4 The thermal conductivity of GA–PMMA nanocomposites as a
function of graphene volume fraction. Inset is the set-up scheme for the
comparative infrared microscopy technique, where the amorphous quartz
having thermal conductivity of 1.3 W/m·K is used as the reference material. To
extract the thermal conductivity of rGO and TBR between rGO and PMMA, the
experimental data is fitted with a modified effective medium theory
simultaneously by a nonlinear least-squares algorithm. The red dashed line
represents the best fitting with the thermal conductivity of rGO (30.0 W/ m·K)
and the rGO-PMMA TBR (3.78×10-8
Km2/W). ............................................ 105
Figure 6.5 The microhardness of GA–PMMA nanocomposites as a function of
graphene volume fractions. Inset is a typical indentation for the Vickers
microhardness test. For all GA–PMMA nanocomposites, the diagonal
dimensions range from 25 to 60 μm. A prediction of the microhardness of
GA–PMMA nanocomposites, using the modified Halpin–Tsai equation, and a
comparative study of microhardness between this work and previously reported
graphene–PMMA nanocomposites are also included. For comparison purposes,
the values of graphene loading reported in mass fraction (wt. %) are converted
to volume fraction (vol. %). ........................................................................... 108
1
CHAPTER 1: Introduction
This chapter provides a background of aerogels and nanocomposites.
Among the various types of aerogels, graphene aerogels, which are a novel type
of carbon-based aerogels, are of particular interest for this study. The utilization
of graphene aerogels to prepare graphene aerogel-based nanocomposites could
effectively solve the agglomeration problem of carbon fillers in carbon-filled
nanocomposites, and further provides a promising application for graphene
aerogels.
1.1 Aerogels
1.1.1 Background
An aerogel is an open-celled, mesoporous and solid foam which comprises
a network of interconnected nanostructures, having a porosity of over 50%.
Herein, the term “mesoporous” (or “a mesoporous material”) is commonly
defined as a material containing pores in the range of 2–50 nm in diameter [1, 2]
.
In fact, aerogels are well-known for their dramatic low densities (often
varying from 0.00016 to ~0.5 g/cm3), and are even considered the lightest solid
materials (with lowest densities) that have ever been made [1, 3, 4]
. For instance,
an as-prepared silica aerogel is only three times heavier than air [5]
, and can
possibly be made even lighter by evacuating the air from its pores. However, a
typical aerogel usually opposes a density of 0.020 g/cm3 or higher,
approximately 15 times heavier than air, and has 95–99% air (or other gases) in
its volume [3]
.
2
It should be noted that the term “aerogel” does not refer to a specific
substance, but rather to a network structure a substance takes on [1]
. Aerogels
can be made of various substances, including silica [5]
, most of the transition
metal oxides [6, 7]
, several main group metal oxides [6]
, organic polymers [8]
,
metals [9]
, carbon [8, 10]
, carbon nanotubes (CNTs) [11, 12]
, and graphene [13, 14]
.
From the perspective of chemistry, the aerogel of a substance is generally
analogous to the bulk of that identical substance. Nevertheless, because of their
extraordinary low densities and length-scale effects, which stem from their
nanostructures, aerogels would commonly exhibit many remarkably improved
properties superior to the non-aerogel form of the substance (e.g. dramatic
increase in surface areas and catalytic activities), while also possibly suffering
reductions in some other properties (e.g. decreased mechanical strength) [1, 2]
.
1.1.2 Carbon-based aerogels
Carbon-based aerogels possess almost similar properties to other kinds of
aerogels, but are also electrically conductive with a density-dependent
conductivity [8, 10-15]
. Due to this unique electrical property, the carbon-based
aerogels are mostly fabricated for industrial applications such as desalination
[16], solar energy collection
[17], and catalyst support
[18, 19]. Noticeably, their high
surface area and electrical conductivity can in particular meet the requirements
of the electrode materials for electrochemical devices [20]
, as shown in Figure
1.1.
3
Figure 1.1 Application examples of carbon-based aerogels. (a) Representation
of an electric double-layer capacitor (EDLC); (b) Three-dimensional N-doped
graphene aerogel-supported Fe3O4 nanoparticles as efficient cathode catalysts
for the redox reaction [20]
. (Reproduced with permission from Ref. 20,
Copyright (2012) American Chemical Society.)
Carbon aerogels were first prepared from organic materials by Pekala et al.
[8, 10] in the 1980s. With this method, resorcinol (R) and formaldehyde (F) were
employed as precursors, sodium carbonate as a catalyst, and water as a solvent.
An organic gel was formed via poly-condensation of the precursors in a solvent
and then cured in an oven for 3–7 days. The as-obtained organic gels were then
supercritically dried to organic aerogels, followed by a carbonization process at
elevated temperature in an inert environment to eventually create carbon
aerogels. From this approach, a high surface area could usually be obtained
from these as-prepared carbon aerogels, which also in turn affected their
electrical conductivity. However, most of the chemicals involved in
4
synthesizing these carbon aerogels are highly toxic, and a high temperature over
1000 °C is also strictly required for the carbonization process.
As a common type of carbon aerogels, CNT aerogels are developed based
on the interesting properties of CNTs [21-23]
. They are essentially CNTs
randomly accumulated in space and connected by some specific organic binder
[11, 12]. It has been indicated that CNT aerogels can offer high power density as
the electrode materials of supercapacitors. However, the high cost of CNT
fabrication and difficulty of CNT dispersion are still great challenges facing the
commercialization of CNT aerogels.
Graphene is a one-atom-thick planar sheet of sp2-bonded carbon atoms
which are densely packed in a honeycomb crystal lattice, having a high surface
area and excellent lateral mechanical, structural, thermal, and electrical
properties, similar or even superior to those of CNTs [24]
. Therefore, graphene
aerogels (GAs) have been widely studied in recent years, as an alternative to the
CNT aerogels [13, 14]
. Unlike CNT aerogels, the GAs can be synthesized from
low-cost graphite at a temperature lower than 200 °C. The extraordinary
properties of graphene and the straightforward and relatively inexpensive
approach give the GAs great potential as the next-generation conducting
aerogels for various applications [15]
, such as energy storage devices [20]
,
actuators and sensors [25]
, and nanocomposites with multifunctional properties
[26].
1.1.3 Drying techniques of aerogels
The critical point drying (CPD), also known as supercritical drying, refers
5
to a special process for creating aerogels by removing the background liquid
from the hydrogel in a gentle and controlled way. It has been widely used to
preserve the porous structures of samples in the production of aerogels,
biological specimens, and microelectromechanical systems (MEMS) [27]
. The
principle of CPD is illustrated in Figure 1.2.
In contrast to the evaporation in the ambient environment, in which surface
tension of the evaporation liquid causes distortion and shrinkage to the gel
network [28, 29]
, the CPD process can allow the liquid phase to pass through a
“supercritical region” instead of crossing the liquid/gas boundary. In the
supercritical region, distinct liquid and gas phases no longer exist, and thus
there is no surface tension applied to cause the collapse of the gel structure [30]
.
However, in practice, the substances which can be used for CPD are still
very limited. This is because most liquids require rather high temperature and
pressure to reach their critical points, which are dangerous and may also destroy
many organic gels. As liquid carbon dioxide (CO2) can be supercritically
extracted at a low temperature of 31.1 °C and a modest pressure of 72.8 atm, it
has been selected as the most common solvent for CPD [31, 32]
.
It also worth mentioning that, since liquid CO2 is immiscible with water,
an intermediate solvent exchange to ethanol, which can be completely miscible
with both water and liquid CO2, is usually required before the real CPD process
to dry hydrogels. The whole process of solvent exchange is carried out by
soaking the hydrogel samples in ethanol for at least several hours and changing
the ethanol at least three times [32]
. The gel after such solvent exchange is called
“alcogel”, and is ready to be subjected to CPD to yield its aerogel form.
6
Figure 1.2 Principles of three drying methods (evaporation, freeze drying, and
CPD) to remove liquid from hydrogels in a generic phase diagram. (a) Direct
evaporation from liquid phase to gas phase causes collapse to the gel network
due to the effect of surface tension. (b) Freeze drying is a common method to
create aerogels. During this process, the samples are fast frozen and then placed
in a vacuum under raised temperature to allow the ice to sublimate. (c) In the
CPD process, the temperature and pressure of liquid carbon dioxide are raised
above its critical point to form a supercritical fluid. By slowly releasing the
pressure, the hydrogels are dried to aerogels with well-preserved networks. The
CPD process is selected as the drying method in this thesis work. Adapted from
the thesis work of Mateusz B. Bryning [33]
.
According to the generic phase diagram illustrated in Figure 1.2, freeze
drying is also a commonly used method to remove the liquid from hydrogels [34]
.
However, Zhang et al. [14]
reported that the aerogels dried by CPD could exhibit
a larger surface area and superior multifunctional property over those that
underwent freeze drying. Therefore, the CPD process has been selected as the
drying method in this thesis. All the CPD processes were completed in a
Tousimis Autosamdri-815, Series B super-critical dryer, following the
procedures recommended by the manufacturer.
1.1.4 General properties and applications of aerogels
Many aerogels have a combination of impressive material properties that
7
other materials may not simultaneously possess. Specific formulations of
aerogels have extreme properties, such as the lowest bulk density (down to
0.00016 g/cm3)
[4], the largest specific surface area of any non-powder material
(as high as 3200 m2/g)
[35], the lowest mean free path of diffusion and dielectric
constant of any solid material [36]
, and the slowest speed of sound through any
solid material [37]
.
Aerogels are generally fragile. Inorganic aerogels are friable and can be
destroyed by bending or a gently applied force [38]
. Organic polymer aerogels
are slightly stronger than those inorganic aerogels. However, the strength of
most types of aerogels can possibly be improved by condensing them at the
expense of their lightweight and ultralow thermal conductivity [39]
.
The real applications of aerogels depend strongly on their various
exceptional properties and features. Table 1-1 below shows some applications
of aerogels, both general and specific. These applications which have been
either demonstrated or proposed previously, arise from the specific properties
or characteristics of the aerogels. In many cases, a specific application may only
be dependent on one single property even if the aerogels have multifunctional
properties which are appropriate to the given application [18]
.
Table 1-1 Identification of aerogel properties and features, with their
applications [18]
. (Reproduced with permission from Ref. 18. Copyright (1998)
Elsevier.)
Property Features Applications
Thermal
conductivity Best insulating solid
Transparent
High temperature
Lightweight
Architectural and
appliance insulation,
portable coolers, transport
vehicles, pipes, cryogenic,
skylights
Space vehicles and
probes, casting molds
Density/ Lightest synthetic solid Catalyst, absorbents,
8
porosity Homogeneous
High specific surface area
Multiple compositions
sensors, fuel storage, ion
exchange
Target for ICF, X-ray
laser
Optical Low refractive index solid
Transparent
Multiple compositions
Cherenkov detectors, light
weight optics, light guides,
special effect optics
Acoustic Lowest sound speed Impedance matchers for
transducers, range finders,
speakers
Mechanical Elastic
Light weight
Energy absorber,
hypervelocity particle trap
Electrical Lowest dielectric constant
High dielectric strength
High surface area
Dielectrics for ICs,
spacers for vacuum
electrodes, vacuum display
spacers, capacitors
The second objective of this thesis work is the utilization of graphene
aerogels to synthesize novel nanocomposites. The next section will present a
brief introduction to nanocomposites.
1.2 Carbon-based nanocomposites
A nanocomposite refers to a composite material for which at least one
dimension of its one component is less than 100 nm, or a distance in nano scale
exists and repeats constantly between its multi-components. In a specific sense,
a nanocomposite is often taken to mean the combination of a bulk matrix and
one or more nanofillers (e.g. nanoparticles, nanosheets, or nanofibers). Because
of the high aspect ratio of the nanofillers, a nanocomposite can typically exhibit
properties in one order greater than those conventional composites [40]
.
To date, carbon-based nanofillers, especially CNTs and graphene, have
been extensively investigated as the multifunctional nanofillers for fabricating
polymer nanocomposites in both industrial and academic fields [41-43]
. However,
9
due to the hydrophobicity and π–π interaction, a homogeneous dispersion of
these carbon nanofillers in the polymer matrix has always been a crucial issue
during the whole fabrication process [44, 45]
. Although many strategies [46, 47]
,
such as solvent processing, in situ polymerization and melt processing, have
been developed, the properties of these nanocomposites prepared through such
powdery dispersion methods are still far less than expected.
The aerogels presented previously may pave the way to motivate the
effective fabrication of the “multifunctional nanocomposite materials”. In
general, an aerogel can first be prepared and optimized as the nanofiller material,
after which the polymer matrix in a liquid or semi-liquid phase could possibly
be backfilled into its network under capillary force [48]
. For the GAs consisting
of graphene nanosheets in particular, this approach provides a novel route for
property enhancements of the graphene-based polymer nanocomposites.
1.3 Objectives of thesis
This thesis aims to achieve an effective morphology–control of the GAs
synthesized by chemical reduction method, and investigate the effects of the
GA morphology on their multifunctional properties. In order to provide the
thermal properties of the as-prepared GAs, a proper technique for thermal
characterization needs to be established. In addition, graphene nanosheets
commonly form agglomerations in the graphene–polymer nanocomposites,
severely limiting their effectiveness. To overcome such agglomeration
problem, a new route for fabricating the graphene-based nanocomposites
should be proposed. The specific objectives of this thesis are as follows:
10
(a) Apply the mild chemical reduction method to fabricate GAs and backfill
polymer into the network of GAs to prepare GA–polymer nanocomposites
by in situ polymerization;
(b) Investigate the effects of synthesis conditions on the morphology and
multifunctional properties of the GAs;
(c) Develop an improved comparative infrared microscopy technique to
characterize the thermal property of the GAs;
(d) Develop and characterize the multifunctional property of GA–polymer
nanocomposites and evaluate how it solves the agglomeration problem of
graphene–polymer nanocomposites prepared by the traditional powdery
dispersion method.
This thesis work would provide a comprehensive and systematical study of
synthesis parameters on the multifunctional properties of the GAs. In particular,
it would provide the first benchmark data of the thermal property of GAs. The
development of GA–polymer nanocomposites may also solve the
agglomeration problem of graphene in the polymer matrix, which would
significantly enhance the multifunctional properties of graphene–polymer
nanocomposites.
It should be noted that this study is based on the selected synthesis method,
while effects of various different synthesis methods on the property of GAs or
GA–polymer nanocomposites are excluded. This work also mainly focuses on
the fabrication and property optimization of materials, rather than the
development of functional applications by the as-prepared materials. In addition,
only existing theoretical models are used to fit with the experimental data in this
11
thesis, and there are no new models established.
1.4 Organization of thesis
This thesis is organized into the following chapters:
Chapter 1 introduces the motivations of this thesis and provides some
background on aerogels, in particular carbon-based aerogels.
Chapter 2 presents a specific literature review of the GAs and
graphene-based nanocomposites, including their structure, synthesis methods,
and multifunctional properties.
Chapter 3 provides a description of the materials, synthesis approach, and
characterization of the developed GAs and GA–PMMA nanocomposites in this
thesis. The improved comparative infrared microscopy technique for thermal
conductivity measurement is presented in detail.
Chapter 4 comprehensively investigates the effects of synthesis conditions,
reducing agents, and CNT inclusions on the morphologies, electrical property,
and thermal stabilities of the GAs. Additionally, the impact of a thermal
annealing treatment at 450 °C for 5h in an Argon (Ar) environment is also
studied. It has been found that the post-thermal annealing treatment and the
additional CNTs can enhance the electrical conductivities of the GAs up to a
factor of 5.
Chapter 5 investigates the thermal conductivities of the GAs with various
graphene volume fractions from 0.67 to 2.50 vol. %, with and without annealing
treatment, measured using a comparative infrared microscopy technique. This
is the first systematical study of the thermal properties of GAs, and the results
12
elucidate the factors limiting their thermal conductivities. The developed
thermal measurement technique can be applied to other porous material
systems.
Chapter 6 presents the fabrication of GA–PMMA nanocomposites by
backfilling PMMA into the pores of the GAs. Due to the uniform distribution of
graphene nanosheets in the PMMA matrix, as graphene loadings increase from
0.67 to 2.50 vol. %, the nanocomposites exhibit significant increases in
electrical conductivities (0.160–0.859 S/m), microhardness (303.6–462.5 MPa),
and thermal conductivities (0.35–0.70 W/m·K) compared with pure PMMA
and the graphene–PMMA nanocomposites prepared by traditional dispersion
methods.
Chapter 7 summaries the major conclusions of this thesis and gives
suggestions for future research work.
The next chapter (Chapter 2) will present a detailed literature review of the
current progress that has been made with graphene, GAs and
graphene–polymer nanocomposites.
13
CHAPTER 2: Literature Review
This literature review chapter includes separate sections on graphene,
graphene aerogels (GAs) and the graphene–polymer nanocomposites: each
section includes subsections addressing the structures, synthesis methods,
multifunctional properties and potential applications of the materials in
question.
2.1 Graphene
2.1.1 Introduction
Graphene is a two-dimensional (2D) planar sheet which is composed of a
monolayer of hexagonally packed carbon atoms in a honeycomb lattice. As
illustrated in Figure 2.1, it has generally been considered as the basic unit for all
graphitic carbon allotropes, e.g. wrapped up into zero-dimensional (0D)
fullerenes, rolled into one-dimensional (1D) carbon nanotubes (CNTs), or
stacked into three-dimensional (3D) graphite [24]
.
Figure 2.1 Graphene is the building block of all carbon materials, e. g. 0D
fullerenes, 1D CNTs, and 3D graphite [24]
. (Reproduced with permission from
Ref. 24. Copyright (2007) Macmillan Publisher Ltd.)
14
Graphene was first isolated from graphite in 2004, using a “Scotch tape”
(or “peel-off”) method [49]
. Since its discovery, graphene has attracted
tremendous attention from both industrial and academic research, due to its
numerous remarkable properties [49-55]
. Defect-free graphene (or “pristine
graphene”) has demonstrated an optical transmittance of 98% [50]
, a large
specific surface area of 2630 m2/g
[51, 52], high electron mobility of 250,000
cm2/Vs at room temperature
[49, 53], an outstanding Young’s modulus of 1 TPa
[54], and superior thermal conductivity of 5300±480 W/m·K
[55]. These unique
intrinsic properties provide graphene with great potential for various
applications, such as field effect devices [56-58]
, transparent conductive films [59]
,
energy storage devices [60, 61]
, and nanocomposites [43, 45, 47, 62-66]
.
2.1.2 Structure of graphene
As shown in Figure 2.2, the honeycomb lattice structure of graphene is
composed of covalently bonded carbon atoms [52, 67, 68]
, with a C–C bonding of
0.142 nm in length. In this structure, all of these carbon atoms are sp2 hybridized
and arranged in a 1s2sp
1sp
1sp
12pz
1 configuration, such that each of them forms
three σ bonds and shares one π bond with its nearest neighbors [69]
. Theory
predicts that these shared π orbitals significantly contribute to the mechanical
stability of the carbon sheet and the delocalized network of electrons [52, 68-71]
.
Although ‘graphene’ refers to a monolayer of carbon atoms in concept,
common references, especially from the experimental point of view, also exist
15
for bilayer, trilayer or few-layer graphene (3–10 layers)1
[72], where several
graphene monolayers stack following the stable orders in natural graphite as an
alternating (ABAB, Bernal type) or staggered (ABCABC, rhombohedral type)
arrangement [72]
with a layer spacing of 3.35 Å [73]
.
Figure 2.2 Scheme of the crystal structure, Brillouin zone, and dispersion
spectrum of graphene [52]
. (Reproduced with permission from Ref. 52.
Copyright (2010) Wiley–VCH.)
2.1.3 Synthesis methods of graphene
There are three common approaches available to synthesize graphene and
graphene-based nanosheets. The following review provides a detailed
consideration of the experimental procedures, strengths and drawbacks of each
technique.
Table 2-1 A comparison of strengths and drawbacks for the methods commonly
applied to synthesize graphene [68]
. (Reproduced with permission from Ref. 68.
Copyright (2010) Elsevier.)
Advantages Disadvantages
Mechanical
exfoliation
Low cost and easy
No special equipment required
Better contrast
Serendipitous
Uneven films
No large-scale production
CVD Even films
Large scale area
Difficult to control
High temperature process
1 In the later sections of this thesis, ‘few-layer graphene’ refers to thin graphene films from 1 to
10 layers.
16
Chemical oxidation
and reduction
Large scale fabrication
Easy suspension handling
Fragile stability
Partial reduction
2.1.3.1 Mechanical exfoliation
The inter-layer Van der Waals interaction energy within graphite is 2
eV/nm2
[74]. Therefore, only a weak force of approximately 300 nN/μm
2 is
required to exfoliate its multilayers to form graphene [74]
. In 1999, Ruoff et al.
[75, 76] first reported the exfoliation of graphite by a micromechanical exfoliation
method, in which graphite crystals were obtained by manipulating
lithographically patterned arrays of graphite micropillars and repeatedly
exfoliating them with a sharp glass tip. Since then, this work has been
considered the foundation of the micromechanical exfoliation method to
fabricate graphene, and also a milestone in outlining of the potential importance
of graphene for various fundamental studies and applications [52]
.
The micromechanical exfoliation method was subsequently modified with
the micropillars of graphite mounted on an AFM cantilever with a purposely
chosen spring constant, so that the pressure and shearing force for exfoliation
could be carefully controlled [74]
. These top–down approaches in general allow
an arbitrary deposition of graphene, yet they cannot be used for fabricating
monolayer graphene or even few-layer graphene of thickness below 10 nm.
Monolayer graphene was first obtained by Novoselov and Geim in 2004 [49, 77]
.
A piece of Scotch tape was used to repeatedly stick and peel a 1-μm-thick
graphite flake, until very smooth and thin fragments were formed. Through
optical microscopy images, the obtained graphene could be subsequently
transferred to arbitrary substrates by gently pressing the tape.
17
For such exfoliation methods, one great challenge resides in the
identification of graphene sheets from the isolated fragments. As the exfoliation
occurs, a large number of fragments can be formed with various layers of
graphene, ranging from 1 to 100 [78]
. Hence the detection of the graphene
monolayer from all the fragments remains difficult. One common method for
detecting the graphene monolayer is the utilization of an optical microscopy [79,
80]. A few previous studies
[24, 49, 53, 77] have shown that when a multilayer
graphene is deposited on a 300-nm-thick silica (SiO2) substrate, its color will
change from yellow to blue under the optical microcopy as its thickness
decreases. Accordingly, when the color appears to change from darker to lighter
shades of purple, few-layer and/or monolayer graphene is indicated, as
illustrated in Figure 2.3 [68]
. The visibility of the graphene originates from the
well-known interference phenomenon [78]
, and is considered definite evidence
of the presence of graphene monolayers. However, a high-quality interface
image is very sensitive to both the thickness and purity of the SiO2 layer, e.g. a
315 nm instead of 300 nm thick SiO2 layer fails to display any apparent contrast
[81]. Therefore, this optical microscopy technique has been gradually replaced
by the Raman spectroscopy technique in recent years, as the latter can provide
more information in terms of both number of layers and structural arrangement
[82]. Another shortcoming of the mechanical exfoliation methods, like the
Scotch tape method, is the presence of glue residue on the graphene sample [83]
,
which not only contaminates the graphene, but also limits its carrier mobility
and other properties [84]
.
18
Figure 2.3 Mechanically exfoliated (a) multi-layer, (b) few-layer and
monolayer graphene. Yellow colors indicate multi-layer graphene of thickness
over 100 nm, while blue, darker, and lighter purple indicate few-layer and
monolayer graphene [68]
. (Reproduced with permission from Ref. 68. Copyright
(2010) Elsevier.)
In general, these mechanical exfoliation approaches may still be the best
ways to generate graphene of the highest electrical and structural qualities
required for fundamental studies. However, due to their time-consuming
process and low yield production, mechanical exfoliation methods are not
highly desired for a scale-up fabrication in practice [68]
. To enhance the ease of
graphene separation from graphite, a liquid-phase exfoliation method has been
developed based on the mechanical exfoliation. During this process, graphite is
dispersed in organic solvents to form suspensions, and the interlayer exfoliation
of graphite is then achieved in the organic suspensions via sonication and/or
centrifugation [85]
.
In 2008, Hernandez et al. [86]
first reported the separation of graphene
layers using N–poly–methylpirrolidon (NMP), in which graphite powders
sifted through a fine sieve were used as the initial material. These graphite
powders were dispersed in NMP to form a homogeneous mixture via sonication.
The mixture exhibited a grey-colored appearance and contained a number of
particles in macroscopic sizes. These particles were subsequently removed from
19
the mixture by centrifugation. A homogenous dark-colored suspension with
only monolayer and few-layer graphene was finally obtained. The properties of
such a graphene–NMP suspension could remain stable for at least five months.
Other organic solvents, such as N, N–dimethylacetamide (DMA),
g–buthyrolactone (GBL) and 1, 3–dimethy1–2–imidasolidinon (DMEU) as the
surface active substances (SAS) for CNTs, have also been utilized in the liquid
phase exfoliation process [78]
. In addition, to improve the quality of the
as-prepared graphene, a post-treatment like thermal annealing at 1000 °C
followed by several purification steps [87]
could be performed to further bring
the graphene dispersed in the suspension to predominantly monolayer
graphene.
The utilization of liquid during exfoliation permits the penetration of
atoms or molecules into the interlayer space of graphite, thus leading to an
enhancement of its interlayer distance and a consequent decrease of the
interaction energy between its neighboring layers [78]
. Therefore, compared with
mechanical exfoliation, liquid exfoliation generally lowers the level of
difficulty in separating the graphite layers. Nevertheless, it still suffers the
disadvantages of being time-consuming and low-yielding, and the separation of
graphene from the organic solvents also remains challenging.
2.1.3.2 Chemical vapor deposition (CVD)
To date, chemical vapor deposition (CVD) is one of the most widespread
methods for the synthesis of carbon-based nanomaterials, such as carbon fibers,
CNTs and graphene. May et al. [88]
reported the observation of a ring-like
20
low-energy electron diffraction pattern on a platinum substrate after thermal
annealing, yet at that time they did not ascribe it to a “monolayer of graphite”.
Later, Blakely et al. [89-92]
investigated the detailed thermodynamics of the
so-called “monolayer graphite” or “bilayer graphite” that was, however, formed
on nickel (Ni) (111) crystals, while the topic of “monolayer graphite” was
proposed as a sub-discipline of surface science for the first time.
So far, great progress has been made on the synthesis of graphene via CVD
processes [93-96]
. One successful example is the large-area growth of graphene
on silicon carbide (SiC) [94, 97, 98]
. When the SiC wafer undergoes thermal
treatment at ~1300 °C in a vacuum environment (or ambient pressure), very
thin graphene with only a few layers can consequently be formed over the
entire wafer. The formation of these graphene samples is considered a
combined result of the sublimation of Si atoms from SiC and the reorganization
and graphitization of the leftover carbon-enriched surface [99-101]
. With careful
control of the process, the graphene monolayer could even be obtained
occasionally through this thermal treatment process [97, 98]
. However, the
graphene obtained from this method exhibits only a misleading and loose
stacking of multilayer structures under Raman and scanning tunneling
microscope (STM) [102, 103]
. Therefore, in order to improve the quality of the
resultant graphene, the method of high-temperature graphitization of SiC still
needs to be modified further [100, 104-106]
.
Currently, the most effective CVD approach is based on the direct
decomposition of hydrocarbons on the metallic substrates [22]
. Land et al. [107]
reported that the decomposition of ethylene on Platinum (Pt) (111) substrates
21
could result in the growth of uniformly distributed graphene islands in
nano-sizes, yet only the layers at their lower step edges appeared to be
continuous. To obtain large single-crystalline graphene, a similar CVD process
decomposing ethylene on Ruthenium (Ru) (0001) substrates was then
conducted by Stutter et al. [108]
The formation mechanism of graphene on
Ru(0001) surfaces was proposed to be analogous to that of the catalysis of
CNTs from transition metal particles [109]
. Ethylene works as the carbon source
in the CVD process and is first dissolved in Ru at 1420 K. As the temperature is
lowered to 1100 K, supersaturation of carbon atoms in Ru is then induced and
subsequently triggers the nucleation of graphene. It has been found that the
graphene formed on Ru(0001) substrates exhibits a large single-crystal with a
lateral dimension exceeding 100 μm. However, due to the strong interaction
between graphene and metal [110]
, the isolation of graphene from Ru(0001)
substrates is usually not easy. This in turn significantly restricts the size of
graphene and also limits the application of graphene for various devices.
Efforts have been made to transfer the epitaxial grown graphene onto
arbitrary substrates. Polycrystalline Ni films [111, 112]
and copper (Cu) [113]
are
applied as the catalyst materials, while centimeter-sized continuous graphene
can be obtained. The mechanism for graphene growing on Ni and Cu is similar
to the dissolution/precipitation mechanism in which Ru is applied as the
template [68]
. As for the substrate transfer, a thin layer of polymer is first coated
on the as-grown graphene, and then the entire sample is dissolved in a dilute
hydrochloric (HCl) solution to remove the Ni substrate [111, 112]
. Owing to the
good dissolubility of Ni in HCl, the metal substrate is easily released. Thereafter,
22
the isolated graphene with polymer coating is attached to the target substrate by
careful pressing, and its polymer coating is eventually removed via immersion
in the organic solvents, e.g. acetone or isopropanol.
Although the CVD process has shed some light on the fabrication of
high-quality graphene over a large area, it still remains challenging to produce
graphene on an industrial scale. In addition, CVD growth of graphene requires
rather expensive templates, and this cost is intrinsically unnecessary for most
low-performance devices. As a result, chemical oxidation and reduction
methods have been developed to produce reduced graphene oxide (rGO), which
is widely considered an alternative to graphene and can be used for most
low-performance applications.
2.1.3.3 Chemical oxidation and reduction
The concept of chemical oxidation/reduction methods to synthesize
graphene basically involves the weakening of the Van der Waals force between
the interlayers of graphene upon insertion of functional groups. When graphite
is oxidized to graphite oxide, its sp2 lattices are partially degraded to sp
2–sp
3
bondings, which consequently reduce the stability of π–π stacking. Following
this, the exfoliation of graphite oxide to graphene oxide (GO) is usually carried
out in suspension under sonication, while the reduction of GO to rGO can be
completed with various reducing agents such as hydrazine and L-ascorbic acid
[114]. Therefore, chemical approaches offer a straightforward and effective route
towards the mass production of graphene [115]
.
The oxidation of graphite can be dated back to the time when Brodie
23
determined the mass of a carbon atom [116]
. It was found that when potassium
chlorate was intercalated in concentrated sulphuric acid and nitric acid, graphite
flakes containing hydroxyl and carboxyl carbon sheets could be produced, with
a large amount of heat released. At that time, these carbon sheets were called
“graphitic acid”, which is basically equivalent to “graphite oxide” today. At
present, the most famous method to synthesize graphite oxide is the one
developed by Hummers et al. [117]
, in which graphite is dispersed into a mixture
of concentrated sulfuric acid (H2SO4), sodium nitrate (NaNO3) and potassium
permanganate (KMnO4), and the mixture is kept at ~45 °C for several hours.
This method has been widely applied for years to fabricate graphite
intercalation compounds, while the products have been commercialized and are
well known as the “Expandable Graphite” [85, 118]
. The volume of the
expandable graphite is over 100 times larger than that of raw graphite, and thus
it becomes a common precursor for the fabrication of GO or rGO.
The exfoliation of graphite oxide to GO could be achieved by various
approaches, such as thermal treatments, sonication or microwaves [119]
. As early
as 1962, a rapid thermal annealing at 1050 °C was applied to split graphite
oxide, and individual sheets were obtained and observed under TEM [120]
.
However, due to its facile processing, sonication has become the current most
popular method, in which the interlayer separation of graphite oxide can be
efficiently fulfilled by sonicating a 0.1–0.5 wt. % suspension for about 12
hours.
The resultant GO aqueous suspensions are usually in yellow and can
remain stable for at least one week. The stability of the GO suspensions can be
24
ascribed to the effect of electrostatic repulsion. In order to further stabilize the
suspensions, some cationic polyelectrolytes such as polyaniline [121]
,
single-stranded DNA [122]
, or polystyrene sulfonate [114]
, and some surfactants
such as polystyrene [123]
, may also be employed. The GO solids can finally be
collected as thin films by deposition [50, 124]
, or as powders by freeze drying.
However, as large numbers of sp3 carbon are present in GO, it intrinsically
exhibits different properties compared with those of the graphene obtained by
mechanical exfoliation or CVD growth. To recover its electrical conductivity
and other properties, a reduction process, either via chemical [114]
or thermal [125]
approaches, is therefore required. Nevertheless, these reduction approaches can
only result in a partial recovery of sp2 configuration, and the rGO products tend
to agglomerate if the reduction is processed in liquids.
2.1.4 Properties of graphene
In recent years, there has been much interest in the development of
graphene and graphene-based materials. This is mainly due to the unique
multifunctional properties that graphene exhibits.
2.1.4.1 Electrical properties
Experimental measurements have shown that, for any graphene obtained
by micromechanical exfoliation, its electron mobility can exceed 2000 cm2/Vs
at room temperature. Similar to CNTs, the electrical conductance in graphene
can also be ascribed to a ballistic transport mode on a scale of ~16 μm [126]
.
Within this length, the electrons in graphene are able to travel without scattering,
25
which leads to its exceptional high change carrier mobility. Otherwise, the
mobility shows a weak dependence when temperature ranges from 10 to 100K,
implying a scattering mechanism solely dominated by its structural defects [127]
.
As the impurities in graphene are minimized by current annealing,
mobility over 25,000 cm2/Vs can be obtained
[128], while with the improvement
of screening technique and skills of substrate removal, a further enhancement is
also possible. Bolotin et al. [84, 129]
have reported a mobility value exceeding
200,000 cm2/Vs for suspended and annealed graphene, which, however, was the
intrinsic limit caused by the scattering of acoustic photons at room temperature
[83, 130]. This remarkable high value is noted to be 10 times greater than that of
copper [131]
. When this mobility is converted to resistivity, a value on the order
of 10-6
Ω·cm is accordingly extracted, which has been considered the lowest
room temperature resistivity ever known. However, the epitaxial grown
graphene typically exhibits a slightly lower electron mobility, which is on the
order of thousands of cm2/Vs
[132], while for the chemically derived graphene
products, their electron mobility only lies on the order of several hundred
cm2/Vs
[133].
2.1.4.2 Mechanical property
As graphene and carbon nanotubes have intrinsically similar lattice
structures, their mechanical properties, such as Young’s modulus and strength,
are generally comparable. The mechanical properties of monolayer and
few-layer graphene have been extensively reported, from both experiments and
modelling.
26
The current widespread values for the Young’s modulus (~1.0 TPa) and
fracture strength (~130 GPa) of the free-standing monolayer graphene were first
measured by Lee et al. [54]
in 2008. A combined technique of nanoindentation
and AFM was applied as illustrated in Figure 2.4 below. From their tests, the
measured stiffness and breaking strength of graphene were 300–400 and ~42
N/m, respectively. After being converted by division with the thickness of the
graphene monolayer, the values were found to be very close to the theoretical
values predicted for bulk graphite [54, 134]
.
Figure 2.4 (a) A SEM image of a monolayer graphene suspended on an array of
circular holes, scale bar: 3μm; (b) A scheme of nanoindentation on the
membrane of graphene [54]
. (Reproduced with permission from Ref. 54.
Copyright (2008) AAAS.)
For chemically derived graphene, i.e. rGO, which is reduced by
hydrogen plasma, the Young’s modulus can only reach one fourth that of
mechanically exfoliated graphene, ~0.25±0.15 TPa [135]
, while its breaking
strength has not been documented yet. Despite their distinctly different
electrical properties, rGO does not show significantly greater mechanical lag
(where mechanical properties are concerned) than the graphene obtained by
mechanical exfoliation. As a result, given its low cost and high mass production,
rGO can be a promising candidate for a reinforcement material [136, 137]
.
27
2.1.4.3 Thermal property of graphene
Base on the estimation by the Wiedemann-Franz law, the contribution
from electrons to the thermal transport in graphene is almost negligible [138]
. The
thermal conductivity of graphene is therefore dominated by the transport of
phonons, which conduct by diffusion at high temperatures and ballistic mode at
low temperatures.
For the pristine graphene, its thermal conductivity (k) has been found to be
proportional to temperature (T) when T is beyond 100 K. Molecular dynamics
(MD) simulations predict a thermal conductivity of 6,000 W/m·K for the
suspended monolayer graphene [21]
, which is dramatically higher than that of
graphitic carbon. Besides the effect of temperatures, many other intrinsic
factors of graphene, such as its dimensions, edge roughness [139]
and edge
shapes [140]
, also significantly impact its thermal property. It has been found that
k follows a power law relationship with the length L, and the exponent lies in the
range of 0.3–0.5 at room temperature.
So far, the most widely accepted thermal conductivity of the pristine
graphene (5,300 W/m·K) is obtained based on the shift of G band from its
Raman spectrum, where the frequency of G peak is measured as a function of
excitation power and the thermal conductivity of graphene is extracted from the
slope of its trend line [141]
. Substrate attachment also plays a significant role in
the thermal conductivity of mechanically exfoliated graphene. When it is
deposited on SiO2 substrates, its thermal conductivity rapidly goes down to
~600 W/m·K, due to the severe phonons leaking and scattering at
graphene-substrate interfaces [141]
.
28
For the graphene synthesized by CVD, a thermal conductivity of ~2500
W/m·K has been reported [142]
, while for chemically derived graphene, there
have been no values reported beyond 10 W/m·K [143, 144]
. In accordance with
these results, the thermal conductivity of graphene is largely determined by the
abundance of sp2 bonds. Therefore, post treatments, such as thermal annealing
[145, 146], plasma annealing
[147], or chemical purification
[87], are effective
approaches for the thermal property enhancement of graphene.
2.1.5 Potential applications of graphene
Because of its outstanding and multifunctional properties, graphene has
become an idea candidate for various applications.
Firstly, as graphene possesses a unique band structure, its electrons and
holes can be highly tunable by a gate electrical field [49]
. Several works have
reported the development of graphene-based field effect transistors (FETs) with
a single black gate [49, 53, 148]
. Levendorf et al. [149]
have proposed the scaled-up
fabrication of graphene transistors based on the epitaxial grown graphene on Cu.
In addition, Lin et al. [150]
have reported the development of high-frequency
(~26 GHz) FETs using top gate geometry. When a polymeric buffer layer was
added between the epitaxial graphene and conventional gate dielectrics, FETs
with cut-off frequencies at ~100 GHz were successfully obtained.
Secondly, due to its high electrical conductivity and extraordinary optical
transmittance in the range of visible wavelengths, graphene also shows great
potential as an important precursor material for fabricating transparent
conductive films (TCFs). Aimed at large scale fabrication, reduced graphene
29
oxide is commonly applied instead of mechanically exfoliated or epitaxial
grown graphene. In addition, various techniques such as vacuum infiltration,
spray coating, spin coating, dip coating etc. have also been employed. Li et al.
[87] reported the preparation of TCFs by Langmuir–Blodgett (LB)–assembling
rGO. The resultant films exhibited a sheet resistance of 8 kΩ and 83%
transparency at 1000 nm wavelength. De et al. [59]
prepared TCFs by directly
sonicating the raw graphite in organic solvents and depositing the dispersions
by vacuum infiltration. The as-prepared films showed a sheet resistance of ~3
kΩ and ~75% transmittance at 550 nm wavelength.
Thirdly, graphene is also a promising candidate for energy storage devices,
due to its large surface area (2630 m2/g) and excellent electrical properties
[60, 61].
So far, there have been many works covering the use of graphene for energy
storage, especially as the electrode materials for electrochemical double layer
capacitors (EDLCs) [151]
. For the electrodes made of rGO reduced by
microwave irradiation or thermal treatment [119]
, high capacitances of up to 190
and 120 F/g have been achieved in aqueous and organic electrolytes
respectively.
In addition, graphene and graphene-based materials have also attracted
extensive attention as fillers in developing graphene–polymer nanocomposites,
which will be reviewed separately in Section 2.4 with more details.
2.2 Graphene aerogels (GAs)
2.2.1 Introduction
As early as 2002, a self-assembly process of three-dimensional (3D)
30
carbon foams was first proposed based on the nanostructured graphite [152]
.
Later, as knowledge of graphene increased, various easier methods to expand
the interlayers of graphite were developed [85, 118]
. Therefore, the direct
preparation of 3D carbon foams using graphene or graphene-based nanosheets
has become attractive in recent years. At present, the reported methods for
fabricating GAs can in general be classified into two categories: with the use of
binders and without the use of binders. Common binders are polymers, such as
polyvinyl alcohol (PVA) [11]
, resorcinol–formaldehyde (RF) [153, 154]
,
ferrocene–grafted poly (p–phenyleneethynylene) (Fc–PPE) [12]
,
and poly
(3,4–ethylenedioxythiophene) (PEDOT) [155]
. The binderless methods can also
be further classified into several sub-categories, like hydrothermal [156]
and
chemical reduction [145, 157]
. However, regardless of which synthesis methods
are used, critical point drying (CPD) or freeze drying must be applied to
preserve the gel structures. After being dried, the resultant GAs combine the
multiple excellent properties of graphene with the high porosity of aerogels,
thus offering a number of outstanding properties, such as extremely low density,
large surface area, super hydrophobicity, and good mechanical properties. So
far, they have been successfully applied to various applications, such as
electronics [158]
, energy storage devices [159, 160]
, catalysis [161]
, and composite
materials [48]
.
Due to its ease of fabrication and dispersion capability, GO (as described
in Section 2.1.3.3) has become the most popular reactant in the graphene family
(pristine graphene, epitaxial grown graphene, etc.) for synthesizing GAs. The
next section provides a brief introduction to the GO, followed by a detailed
31
literature review of work concerning GAs.
2.2.2 Graphene oxide (GO)
As described in Section 2.1.3.3, graphene oxide (GO), the intermediate
material of rGO before reduction, is considered oxidized graphene having
various reactive oxygen functional groups such as carbonyl, hydroxyl, and
epoxide groups on its surface. It is normally obtained from the oxidative
treatments of graphite and successive exfoliation. As expected, the oxygen
content of the GO, or the C: H: O compositions, exhibits slightly different
through varying procedures. However, a C: O ratio of ~2:1 has been found to be
the upper limit of the overall extent of oxidation, regardless of one-stop or
multiple oxidation approaches [85, 116-118]
.
Due to its complexity and variability (even from sample to sample), the
structural features of GO have remained debatable for many years, and even
until today. The current most well-known model was proposed by Lerf and
Klinowski [162, 163]
, and is shown in Figure 2.5 below. Two sub-models, one
incorporating the existence of carboxylic acid groups and one not, are presented.
The dominant sub-model largely depends on the operative oxidizing conditions.
Apart from the Lerf-Klinowski model, other models such as the Hofmann [117]
,
Ruess [164]
, and Nakajima-Matsuo [165, 166]
models, and modifications [167, 168]
of
the Lerf-Klinowski model, have also been proposed to explain the structure of
GO. However, the differences in raw materials and reaction pathways upon
oxidation cause substantial variance in their exact structures and properties, and
thus difficulties with the identification of a unique structural feature of the GO
32
remain [169]
.
Figure 2.5 Two sub-models of the Lerf-Klinowski model with (top) and
without (bottom) the presence of carboxylic acid groups on the periphery of the
graphitic basal plane of GO [162, 163, 169]
. (Top: Reproduced with permission from
Ref. 162. Copyright (1998) American Chemical Society. Bottom: Reproduced
with permission from Ref. 163. Copyright (1998) Elsevier.)
To date, there have been many approaches proposed for the preparation of
GO. Among them, the most popular one is the modified Hummers’ method [117]
,
which applies a combination of potassium permanganate (KMnO4) and
concentrated sulfuric acid (H2SO4) to oxidize graphite. Operative procedures of
the Hummers’ method will be presented step by step in Section 3.2.1 and thus
not described here. Notably, the starting material commonly used to synthesize
GO is the graphite flake, which contains numbers of localized defects in its π
lattice serving as the seed for oxidation [170]
. Due to the abundance of
oxygen-containing functional groups, GO possesses good hydrophilicity in
water and aqueous liquids [162, 171]
. It has been deduced that the strong bonding
33
between water (H2O) and the basal lattice of GO is dominated by the oxygen in
its epoxides, as illustrated in Figure 2.6.
Figure 2.6 Schematic of hydrogen bonding formed between water (H2O) and
the oxygen-containing functional groups of GO [169]
. (Reproduced with
permission from Ref. 169. Copyright (2010) The Royal Society of Chemistry.)
Although these functional groups provide GO with a good hydrophobicity,
they also disrupt the sp2 network of the graphitic lattice, consequently resulting
the GO in an electrically insulating material with a conductivity of
0.0206±0.002 S/m [125]
. As presented in Section 2.1.3.3, a reduction process,
either chemical [114]
or thermal [125]
, has been found to be conducive to the
partial recovery of the conductivity of GO . However, as the oxidation process
itself is detrimental to the graphitic lattice planar [125, 172]
, the multifunctional
properties of rGO are generally similar to, but expectedly fall behind, those of
the pristine graphene.
In terms of the synthesis of GAs, GO obtained through chemical
conversion is still the most commonly used starting material, owing to its facile
preparation, excellent hydrophobicity, and other good properties. In addition,
its chemical and thermal reduction routes also pave the way to different
34
synthesis methods of the GAs, which will be presented in detail in Section 2.2.4.
Above all, the presence of functional groups on GO facilitates its multiple
chemical transformations to synthesize graphene-like materials including GAs,
thus offering great potential for various large scale applications.
2.2.3 Structure of GAs
Similar to most other aerogels such as silica aerogels, carbon aerogels and
CNT aerogels, GAs are also constructed by their substance, i.e. graphene, but in
the form of interconnected networks. As shown in Figure 2.7(a), the GAs are
intrinsically highly porous, and graphene nanosheets are randomly distributed
within their frames. So far, the graphene type used to build the blocks of GAs is
mostly rGO, rather than mechanically exfoliated graphene or epitaxial grown
graphene. This is not only due to the low cost and high production rates of rGO,
but also because rGO nanosheets themselves are able to self-assemble upon
reduction [114]
. Figure 2.7(b) shows the thickness of an rGO monolayer, close to
1nm. Although this value is nearly 3 times higher than the theoretical thickness
of the pristine graphene, such a nanosheet is still widely acceptable as the basic
unit of GAs, given the presence of functional groups, such as hydroxyl,
carboxyl and epoxide groups, on its surface [157]
.
The morphologies of the GAs are significantly affected by their synthesis
routes, including the method selected and its reaction conditions [173, 174]
. The
GAs usually have surface areas in the range of 100–1200 m2/g
[174, 175], which
are lower than that of monolayer graphene (~2630 m2/g), indicating that overlap
of varying severity occurs with the graphene nanosheets in GAs [61]
. A
35
systematic study of the relation between the morphologies and reaction
conditions of the GAs has been conducted for this thesis, and will be presented
in Chapter 4.
Figure 2.7 (a) A typical SEM image which shows the structure of GAs. (b)
AFM images of typical building blocks of the GAs [157]
. (Reproduced with
permission from Ref. 157. Copyright (2011) Elsevier.)
2.2.4 Synthesis methods of GAs
2.2.4.1 Crosslinking induced method
Inspired by the synthesis of organic (resorcinol (R)–formaldehyde (F))
aerogels, Worsley et al. [153, 175]
reported the synthesis of ultralow-density GAs
using R and F as an organic binder that is to produce carbon cross-links between
graphene nanosheets. In this method, the gelation process exactly follows the
procedures to prepare R–F organic aerogels [8]
, except that GO is first dispersed
in the aqueous solutions of R and F. After the processes of solvent exchange and
supercritical drying (or freeze drying), the resultant dry gels are basically
GO–RF aerogels. A pyrolysis at 1050 °C under inert gases was subsequently
required to carbonize the GO–RF aerogels into carbon-based GAs. Besides RF,
other polymers such as PVA/poly (styrenesulfonate) (PSS) [122]
and pluronic
copolymers [176]
have also been reported as the binders to stabilize GO and
36
fabricate GAs. However, as most of these binders used for linking graphene
nanosheets are toxic polymers and an extremely high temperature, over 1000
°C, is required for their carbonization, this method has been largely replaced by
several binderless methods.
Tang et al. [177]
and Jiang et al. [178]
developed an ion linkage method to
induce crosslinks between the graphene nanosheets. Under a hydrothermal
process, noble particles are present in situ and facilitate the nucleation and
assembly of the graphene nanosheets. Within the GAs, these particles also play
a role to decorate the GAs, rather than only serving as a binder. Therefore, these
embedded particles can give the resultant GAs further potential for various
functional applications, such as sensors [179]
and catalysis [180]
.
2.2.4.2 Hydrothermal method
In 2011, Xu et al. [156]
reported a one-stop hydrothermal technique to
prepare binderless GAs. A GO aqueous suspension of 2 mg/ml was sealed in a
Teflon-lined autoclave and heated at 180 °C for 12 h. After this hydrothermal
process, graphene hydrogels were formed as a result of hydrothermal reduction
and by physical links via π–π interaction. Finally, these hydrogels needed to
undergo solvent exchanges and either supercritical drying or freeze drying to
yield the GAs.
Based on this hydrothermal method, Nguyen et al. [174]
conducted a similar
synthesis process to investigate the effects of GO concentrations and
hydrothermal treatment time on the morphologies of the as-prepared GAs. It
was found that the surface areas and total pore volumes of the GAs increased
37
with the increase in GO concentration, but decreased with the increase in
treatment time. However, when the GO concentration was lower than 0.5
mg/ml, GAs could not be formed during such a hydrothermal process [156]
.
Since its development, the hydrothermal method has been considered a very
effective approach to synthesize binderless GAs. However, its requirements for
high temperatures and high pressure during the operation may still strongly
limit its potential to be scaled up.
2.2.4.3 Chemical reduction method
Zhang et al. [14, 157]
reported an easy and environmental-friendly method to
synthesize GAs. The operation of this synthesis process is quite simple. GO
nanosheets were first dispersed in DI water to form aqueous suspensions, and a
reducing agent such as LAA was then added to the suspension. After being kept
still for several hours either at room temperature or low temperatures below 100
°C, graphene hydrogels were then formed via the self-assembly of graphene
nanosheets. After that, solvent exchanges and supercritical drying (or freeze
drying) were also needed to dry these hydrogels to aerogels.
As with the hydrothermal method presented above, the reaction conditions
of this chemical reduction process also significantly impact the morphologies
and properties of the GAs. Chen et al. [145]
investigated the effects of various
reducing agents on the gelation time required to form the graphene hydrogels.
They found that the gelation of L-ascorbic acid (LAA)– and sodium sulfide
(Na2S)–reduced samples took only 10 min, while the sodium hydrogen sulfite
(NaHSO3)-reduced ones took 30 min. The lowest GO concentration needed to
38
form hydrogels was reported by Zhang et al. [181]
, where the gelation of rGO
nanosheets was successfully achieved via the reduction of the GO concentration
of as low as 0.1 mg/ml using oxalic acid and sodium iodide (NaI).
This chemical reduction method is considered a mild and ‘green’ way to
fabricate the GAs. However, at present, very few studies have been conducted
on the effects of synthesis parameters on the morphology and multifunctional
properties of the resultant GAs. The impact of different reducing agents on the
GA nanostructures has also yet to be investigated.
2.2.4.4 Other synthesis methods
Besides the three methods described above, there are also several other
methods to synthesize the 3D structures of GAs. When the GO concentration is
high enough (e.g. 30 mg/ml), it can self-gelate to hydrogels via partial π–π
stacking with the aid of a prolonged sonication [182]
. However, due to the
abundance of functional groups on the GO, it is usually difficult to achieve high
electrical conductivities and mechanical strengths for these aerogels after
supercritical/freeze drying.
A free-standing monolith of the graphene network can also be obtained via
a CVD process on various templates such as Ni or Cu foams [183]
, nanosilica [184]
,
and zinc oxide (ZnO) tetrapods [185]
, which basically follows the CVD of
hydrocarbons to produce graphene. After graphene is successfully formed on
these templates, acids can then be applied to remove the corresponding
framework materials, consequently leaving the 3D architectures of graphene.
Although this CVD method can yield graphene of relatively high quality,
39
suitable templates for the growth of the desired graphene networks are usually
difficult to find.
Another template-direct method can be achieved by directly freeze drying
the GO suspensions [186, 187]
, which is also known as the “ice template” method.
He et al. [187]
reported a systematic study to investigate the effect of freezing
routes on the structures and properties of the aerogels. Interestingly, they even
obtained the GAs with highly ordered pores, under a unidirectional freezing
condition.
2.2.5 Properties of GAs
The multifunctional properties of GAs strongly depend on their synthesis
methods and structures. The following sections present a major review of the
electrical, mechanical and thermal properties of the GAs.
2.2.5.1 Electrical property of GAs
Due to different synthesis methods and measurement techniques, the
electrical conductivities of GAs are reported to be within a large range of
0.1–100 S/m [145, 156, 175, 188]
. For the GAs synthesized with the R-F method, a
remarkable electrical conductivity of up to 87 S/m [188]
is obtained by a
four-probe technique, while the electrical conductivity is only ~0.1 S/m [156]
for
the hydrothermally synthesized samples measured by a two-probe technique.
Such large variations may be partly due to the characterization methods, as
contact resistance may probably exist in the two-probe technique. Apart from
this reason, the different bondings formed during different synthesis approaches
40
of GAs should be intrinsically responsible for their conductivities [15]
.
Efforts [145, 189]
have also been made to investigate the effects of drying
methods and reducing agents on the electrical properties of GAs. It has been
found that supercritical drying could yield a more conductive sample compared
with freeze drying [189]
, while the utilization of a more efficient reducing agent
would also ensure the GAs with a higher electrical conductivity [145]
. To further
enhance the electron transfer in GAs, post treatments such as thermal annealing
under inert gases could be performed on the as-prepared GAs. It has been
reported that the functional groups on GO could be effectively removed by
thermal treatments above 200 °C [190]
. Therefore, annealing the GAs in an inert
environment could further purify the graphene sheets and increase their
electrical conductivity. Chen et al. [145]
reported that a thermal annealing
process at 400 °C enhanced the electrical conductivity of the GAs up to nearly 5
times.
2.2.5.2 Mechanical property of the GAs
Since GAs are constructed from individual pieces of rGO, their
mechanical properties, including Young’s modulus and strength, are far less
than that of the graphene monolayer [54]
. Therefore, the GAs are usually
addressed to be fragile and brittle, and polymer binders may be needed to
reinforce their networks. Nevertheless, given the extremely low densities of the
GAs, their specific strength would become impressive. For example, the GAs
reduced by LAA and dried by supercritical CO2 have been reported to be able to
support over 14,000 times their own weights with little deformation [14]
. In this
41
same work [14]
, different drying methods were also performed and compared to
study their effects on the mechanical performance of GAs. It has been found
that those GAs processed by freeze drying can only support 3,300 times their
own weight, 4 times less than those dried by supercritical CO2 [14]
.
Due to the ease of operation, most mechanical properties of the GAs are
obtained from their compressive tests. The compressive stress-strain curves of
GAs can usually be divided into 3 regions: the elastic region, yield region and
densification region [145]
. The elastic region lies at low compressive strains
where the GAs elastically behave. The region from the end of the elastic region
to ~60% compressive strains is usually considered the yield region, where the
nano-sized pores in the GAs start to collapse and the increase in stresses appears
to slowly accelerate. Finally, the densification region is defined as the region
from ~60% compressive strains onwards, where almost all the pores have
collapsed and the GAs are completely hardened by the further compression.
From the compressive tests, the measured Young’s modulus and strength of
GAs normally vary from 0.26–20 MPa and 0.042–2.2 MPa, respectively [25, 156,
191].
2.2.5.3 Thermal property of the GAs
GAs with continuous scaffolds and mesoporous structures may reduce the
internal thermal resistance and enhance the thermal efficiency of TIMs.
However, studies of the thermal transport in GAs still remain very limited.
Zhong et al. [192]
reported the thermal conductivity of a GA sample with ~11
vol. % graphene and a relatively low surface area (~43 m2/g) to be 2.18 W/m·K
42
using the laser flash technique.
To characterize the thermal conductivities of nanostructured materials [193]
,
several measurement techniques have been developed, including the laser flash
technique [192, 194, 195]
, steady-state measurement techniques [196, 197]
, transient
techniques [198, 199]
, and infrared (IR) microscopy techniques [62, 200, 201]
. IR
microscopy has several advantages over the other methods as it uses the
non-contact, two-dimensional temperature mapping, eliminating the need for
intrusive temperature sensors, and is a direct measurement of thermal
conductivity (e.g. it does not require information about the specific heat and
density). IR microscopy techniques have been utilized to measure the effective
thermal conductivities of a bulk material [62]
and thermal resistances of
commercial TIMs [201]
with a reported approximated uncertainty of 10%. For
thermal conductivity measurements in particular, the heat flux can be extracted
even more accurately with the employment of reference materials based on a
comparative method similar to the ASTM E1225 standard. However, the
utilization of IR microscopy techniques to measure the thermal conductivities
of GAs has not been reported yet.
2.2.6 Potential applications of GAs
Because of their large surface areas and good electrical conductivities,
GAs have attracted great interest in the area of energy storage, especially as the
electrode materials for supercapacitors. Meng et al. [19]
prepared alkali-treated
GA electrodes in supercapacitor cells, and their specific capacitance was
measured to be 122 F/g at a discharge currency density of 50 mA/g. Sun et al.
43
[202] tested the electrochemical performance of GAs in a propylene carbonate
electrolyte and obtained a specific capacitance of 140 F/g at a current density of
1 A/g. In addition, Zhang et al. [203]
assembled the supercapacitors using a GA
synthesized via hydrothermal process and reported an even higher specific
capacitance of 220 F/g at 1 A/g. These supercapacitors were also found to be
highly stable, and still retained 92% capacitance after 2000 cycles.
GAs can also be used as a thermal enhancer and container for phase change
materials. The GA–octadecanoic acid (OA) composites have been found to
effectively enhance the thermal conductivity of OA by ~14 times, and their
latent heat is close to the value of pure OA [192]
. Many other research works have
used GAs as absorbents [204]
, stain sensors [205]
, and fire-resistant materials [13]
.
Up to now, efforts are still being made to improve the multifunctional properties
of GAs, which will promisingly push the GAs forward to more practical
applications in the future.
In addition, the emergence of the GAs may also provide a new type of
reinforcement material to solve the aggregation problems in composites. A
uniform distribution of graphene fillers in the composites can probably be
achieved by backfilling polymer matrices into the pores of the GAs.
Nevertheless, the development of GA-based nanocomposites has not been
widely reported so far. A comprehensive study of the GAs and the development
and optimization of GA–PMMA nanocomposites would therefore be essential.
44
2.3 Graphene–carbon nanotube (CNT) hybrid aerogels
2.3.1 Introduction
As both graphene and CNTs have similar extraordinary electrical,
mechanical, and thermal properties [206, 207]
, efforts have also been made to
combine them to achieve a synergistic effect [208, 209]
. So far, graphene–CNT
hybrid aerogels, 3D networks of graphene combined with CNTs, have been
successfully developed by many groups worldwide, showing good performance
in various applications, such as for lithium ion batteries [210]
, transparent
conductors [211]
, and capacitive deionization (CDI) electrodes [157, 212]
for water
purification.
2.3.2 Structure of graphene–CNT hybrid aerogels
The networks of GAs are formed by the self-assembly of graphene
nanosheets due to hydrophobicity and π–π interaction. Herein, for these
graphene–CNT hybrid aerogels, as shown in Figure 2.8 below, CNTs are most
physically wrapped in the GA networks, without inducing any structural change.
However, their existence may possibly prevent the stacking of graphene
nanosheets and, remarkably, enhance the surface areas of the GAs. Recently,
the formation of in situ chemical bonding between graphene and CNTs has been
reported by Yan et al. [209]
. A carefully controlled hydrothermal process has
been found to result in a removal of most functional groups and defects in both
the graphene nanosheets and functionalized CNTs. Therefore, these two carbon
allotropes can possibly be bonded after the simultaneous reduction.
45
Figure 2.8 SEM images showing the network structures of graphene–CNT
hybrid aerogels [157]
. (Reproduced with permission from Ref. 157. Copyright
(2011) Elsevier.)
2.3.3 Synthesis methods of graphene–CNT hybrid aerogels
The synthesis of graphene–CNT hybrid aerogels can basically follow all
the approaches to synthesize GAs [153, 156, 157]
. Among them, the chemical
reduction method is usually selected for its facile operation and high-quality
production [157, 212]
. For a typical procedure, CNTs are first dispersed in a GO
aqueous suspension by sonication, while the mixture is then reduced by
reducing agents, such as L-ascorbic acid (LAA) or hydrazine hydrate. It has
been found that CNTs can be well dispersed in the precursors with the presence
of GO nanosheets. Therefore, the existence of a synergistic effect between
CNTs and graphene is accordingly proposed. Besides the chemical reduction
method, there some works also report the synthesis of graphene–CNT aerogels
by directly freeze drying the CNT–GO mixture [4]
, or using a hydrothermal [209]
or two-step CVD method2
[213]. Further details of these methods can be found in
Section 2.2.4.
2 In fact, the CVD approach can only lead to the formation of macroporous structures, instead
of aerogels.
46
2.3.4 Properties and potential applications of graphene–CNT hybrid aerogels
It has been found that the CNTs incorporated into GAs can serve as a
bridge between the graphene nanosheets [210]
. Therefore, such hybrid materials
not only get access to the more effective electron transfer, but also minimize the
agglomeration problems of CNTs/graphene, and further increase their surface
areas. As addressed by Yoo et al. [210]
, the incorporation of CNTs with graphene
nanosheets can effectively increase the layer distance between the graphene
sheets (as shown in Figure 2.9 below), thus increasing the specific capacity of
lithium ion batteries up to ~5 times. In addition, the transparent conductive
films (TCFs) developed by combining CNTs with chemically derived graphene
can exhibit an electrical resistance of 240 Ω at 86% transmittance, which is very
impressive and almost comparable to the performance of commercial indium tin
oxide (ITO) [211].
Figure 2.9 Relationship between d spacing of graphene nanosheet (GNS)
families and graphite [210]
. Addition of CNTs and fullerene (C60) was found to
effectively increase the interlayer spacing of graphene. (Reproduced with
permission from Ref. 210. Copyright (2008) American Chemical Society.)
The most recently developed graphene–CNT hybrid aerogels show a large
surface area of 435 m2/g and high conductivity of 7.5 S/m
[212]. When used for
47
water purification, they surprisingly provide a desalination capacity of 633.3
mg/g, with the concentration of sodium chloride (NaCl) reaching 35 g/L, over
15 times higher than that of most well-known CDI materials [16]
. Therefore,
these graphene–CNT hybrid aerogels hold great potential in the field of water
purification, including CDI of light metal salts, removal of organic dyes, and
enrichment of heavy metal ions in water sources [212]
.
2.4 Graphene–polymer nanocomposites
2.4.1 Introduction
As presented in Section 2.1.5, one of the promising applications of
graphene is to incorporate it with polymers to fabricate graphene–polymer
nanocomposites [43, 45, 47, 62-66]
. Compared with CNTs, graphene seems to be a
more favorable candidate due to its multiple fabrication routes and various
derived products, including GO and rGO as alternatives [43]
.
Up to now, various polymers such as PVA [214, 215]
, Poly(methyl
methacrylate) PMMA [216]
, Polystyrene (PS) [217]
, Polypropylene (PP) [218]
, and
epoxy [219]
have been used to fabricate graphene–polymer nanocomposites,
mainly through three approaches [46, 47]
: (1) solvent processing, (2) melt
processing, and (3) in situ polymerization. The multifunctional properties of
these developed nanocomposites have also been extensively investigated and
compared with the nanocomposites fabricated with CNTs. To enhance some
particular properties of the nanocomposites, efforts are still being made toward
the selection of graphene type, modification of the graphene surface, and
improvement of graphene–polymer interaction. These developed
48
graphene–polymer nanocomposites show good performance in many aspects,
such as electrical, electrochemical, and thermal properties, and thus offer great
potential for many applications, including flexible electronics [220]
, energy
storage devices [221]
, and thermal interface materials (TIMs) [201]
.
2.4.2 Structure of graphene–polymer nanocomposites
The combination of the pristine graphene and polymers is normally
facilitated by the crosslinks formed by the polymers themselves. However, how
uniformly the graphene is dispersed in the polymer matrices has remarkable
effects on the structures and properties of the resultant nanocomposites, as
agglomeration of hydrophobic nanofillers in polymer matrices has been
extensively reported [222]
. Although attempts such as surface modification and
radiation treatment have been made to improve the solubility of graphene, a
uniform distribution of these nanosheets is still very difficult to achieve [223]
.
Sonication may temporarily provide a versatile way to improve the dispersion,
but often causes damage to the nanosheets and weakens their intrinsic
properties [223]
.
Currently, the rGO is regarded as a good alternative nanofiller, instead of
the pristine graphene, because the presence of various functional groups such as
epoxide and hydroxyl groups on the rGO can be very effective for its dispersion
[224]. For some cases without crucial requirements for electrical conductance,
hydrophilic GO can also be directly applied without reduction [41]
.
49
2.4.3 Synthesis methods of graphene–polymer nanocomposites
2.4.3.1 Solvent processing method
In solvent processing, graphene-based fillers are usually first dispersed in
an organic solvent by sonication, after which the suspension is mixed with a
polymer solution by magnetic agitation or sonication. The graphene–polymer
nanocomposites are thereafter obtained by evaporating or distilling the organic
solvent in the mixture. Various polymers, such as PP [225]
, Polyvinylidene
fluoride (PVDF) [226]
, and PMMA [227]
, have been reported to be suitable for this
solvent processing. However, the removal of the solvent without disturbance of
graphene distribution is still critical for this method.
2.4.3.2 Melt processing method
Melt processing is a versatile method, which is believed to be
environmentally friendly and appropriate for the mass production of polymeric
materials, especially thermoplastics [224]
. Graphene-based composites using
polymers (polyurethane (PU) [228]
, polyethylene terephthalate (PET) [229]
, Poly
Carbonate (PC) [230]
etc.) have been successfully prepared with this method.
This method applies a high temperature and high shear forces to disperse
graphene nanofillers in the polymer matrix, and avoids using other toxic
solvents. Compared with solvent processing, melt mixing is inferior in
effectively dispersing graphene, as the viscosities of the molten polymers are at
times too high to intercalate the fillers [45]
. However, for the preparation of the
composites with low filler loadings, melt processing is still one of the most
practical approaches, and can be readily scaled up for industrial manufacture.
50
2.4.3.3 In situ polymerization method
In situ polymerization is another efficient way to synthesize
graphene–polymer nanocomposites, as graphene fillers can be uniformly
dispersed and form strong interaction with the polymer matrix [45]
. In this
method, graphene is first mixed with monomers or pre-polymers, while the
composites are subsequently obtained by a complete polymerization [231-233]
.
Composites such as rGO–epoxy [234]
, rGO–PMMA [216]
and rGO–PU [235]
have
been successfully developed via this method, and the dispersion of graphene
nanosheets in their polymer matrices has been found to be improved. However,
the problems arising from the solvents during polymerization, similar to those
of the solvent processing method, still remain challenging. As the viscosity
increases with the processing of polymerization, the graphene loading which
can be uniformly dispersed is actually low.
2.4.4 Properties of graphene–polymer nanocomposites
2.4.4.1 Electrical properties
When applied as nanofillers in insulating polymer matrices, graphene can
significantly increase the electrical conductivity of the composites [46]
.
Considerable research [236-238]
has been conducted to develop graphene-based
nanocomposites for electrical applications. For example, the rGO–PS
nanocomposites are reported to show electrical conductivities up to 1 S/m at 2.5
vol. % rGO loading [236]
, while the rGO–epoxy composites have a conductivity
of 10 S/m with rGO concentration of 8.8 vol. % [239]
. Because of the wide
selection of polymers, fabrication methods and measurement techniques
51
applied, the measured electrical conductivities of the graphene-based
nanocomposites are highly scattered. For rGO–PS composites alone, the values
are reported to vary in the range of 10-4
-72 S/m for the composites within 2.5
vol. % rGO fraction [236, 237]
. Therefore, besides the intrinsic conductivities of
the fillers, the effectiveness of the enhancement is probably also determined by
the distribution of the rGO nanosheets in the polymer composites.
To evaluate the filler distribution, the percolation threshold (φc) is always
considered an important parameter [197, 239, 240]
. It indicates the point at which a
conductive network of the fillers is formed in the composites, and thus suggests
the occurrence from insulation to conduction. For most cases, the lower the
threshold φc, the more uniform the dispersion and the more effective the
enhancement. The lowest threshold was reported by Dao et al. [241]
, which is
0.04 vol. % for rGO in PMMA microsphere.
Recently, it has been found that the reduction process of the GO also
impacts the percolation threshold of the rGO–polymer nanocomposites [242]
.
The chemically reduced GO has a lower threshold than the thermally reduced
GO. As it is well known that thermal reduction can remove more functional
groups of GO and lead to more restacking, the higher threshold of its resultant
composites is probably proof of the increased difficulty in filler dispersion [242]
.
So far, effective enhancement of the electrical properties of graphene-based
nanocomposites still remains a challenge.
2.4.4.2 Mechanical properties
It has been found that the inclusion of graphene or graphene-based
52
nanofillers can significantly influence the strength and elastic modulus of the
composites. Nevertheless, the reported data for such improvements are
scattered. A modest strength increase of only 7 % has been reported for a 3 wt. %
GO–Polyimide (PI) composites [243]
, while a remarkable increase of 212 % has
been reported for a 2 wt. % rGO–PVA [244]
, and an increase as high as 893 % for
a 3 wt. % Octadecylamine (ODA)–GO–PI composite [243]
. Similarly, for their
elastic modulus, enhancements varying from 20 % [245]
to over 1400 % [243]
have
been reported for different composite compositions. These significant variants
are probably due to the distinguished qualities of graphene and the different
polymerization approaches involved in the processing.
The mechanical properties of the GO and rGO are inferior but still
comparable to those of the pristine graphene. The 0.7 wt. % GO–PVA
composites show an increase of 76% and 62% in strength and modulus
respectively, compared with pure PVA [238]
. Bao et al. [214]
compared the 0.8 wt. %
loading of GO and 0.8 wt. % rGO in PVA. It was found that the GO–PVA and
rGO–PVA composites exhibited an increase of 54% and 66% in strength and 52%
and 70% in modulus, respectively. As a result, the researchers reported that the
absence (or minimal presence) of functional groups on the graphene lattice may
help to restrict and order the chain arrangements in polymers, thus leading to
these minor enhancements.
2.4.4.3 Thermal properties
Thermal management is a crucial issue in the electronics industries, owing
to the continued miniaturization and rapid increase in the power of
53
microelectronics, optoelectronics, and photonic devices [246, 247]
. Recently, due
to the remarkable thermal conductivity of graphene (~5300 W/m·K [55]
),
graphene–polymer nanocomposites have been considered an ideal candidate for
next-generation thermal interface materials (TIMs) [201]
.
Much effort has been made to evaluate the thermal performance of the
graphene-based nanocomposites [42, 194, 248, 249]
. Yavari et al. [248]
reported that
the thermal conductivity of graphene–1–octadecanol composites was increased
by ~140% upon ~4 wt. % graphene loading. Yu et al. [249]
applied graphene
nanoplatelets in epoxy and found that the thermal conductivity of the resultant
nanocomposites reached up to 6.44 W/m·K, at ~25 vol. % graphene loading.
However, the enhancement in thermal conductivity for graphene-based
composites is still very limited because of several factors, including local
agglomeration, defects within the graphene nanosheets, and interaction between
the graphene nanosheets and the polymer matrix [141, 183, 222]
.
The thermal boundary resistance (TBR, known as Kapitza resistance [250]
)
between the nanofillers and polymer matrix explains the large discrepancy
between the measured and theoretical thermal conductivity of polymer
composites. However, due to experimental limitations, TBRs between graphene
nanosheets and polymer matrices are generally difficult to measure. The
effective medium theory (EMT) developed by Nan et al. [251]
is one of the most
common approaches to estimate these TBR values, when given the thermal
conductivity of the composite. In Nan’s approach, the dimensions (e.g. length,
width and thickness) of graphene nanosheets are simplified and fixed to be
single input values, which is convenient for application in various composite
54
systems.
2.4.5 Potential applications of graphene–polymer nanocomposites
The applications and potential of graphene–polymer nanocomposites are
strongly dependent on their properties, which have been discussed. Notably,
most applications of these graphene-based composites are similar to those for
CNT-based composites. However, the diverse approaches to obtain graphene
make it more ideal than CNTs for in-practice applications.
Based on their significantly enhanced electrical conductivities, the
development of graphene-based composite electrodes is one of the most widely
investigated applications. For instance, an rGO–PANI electrode is reported to
have a specific capacitance of over 1000 F/g [252]
. Otherwise, since conductive
polymers normally possess a specific temperature coefficient, they can also be
developed as temperature sensors, e.g. rGO–PVDF composites [26]
. In addition,
when the resistivity of a material is lower than 105 Ω/square, it can be used for
EMI shielding. It has been reported that an rGO–epoxy composite with 15 wt. %
graphene loading can reach the same competitive level as the commercial
materials [239]
.
The enhancement of the mechanical properties of polymer matrices gives
graphene-based nanocomposites great potential as advanced structure materials,
which are light-weight with high strength and modulus [243, 244]
. Additionally,
based on the improved thermal properties of graphene–polymer
nanocomposites, efforts are also being made to develop them for thermal
management applications [42, 248, 249]
. As these applications concern the aspects
55
of the mechanical and thermal properties that can be directly referred to in
Section 2.4.4, the details will not be repeated here.
The next chapter will present the experimental section of this thesis.
56
CHAPTER 3: Experimental Section
3.1 Materials
Graphite powder (CAS No. 7782-42-5, < 20 μm), sodium nitrate (NaNO3,
CAS No. 7631-99-4, ≥ 99.0%), potassium permanganate (KMnO4, CAS No.
7722-64-7, ≥ 99%), concentrated sulfuric acid (H2SO4, CAS No. 7664-93-9, ≥
99.999%), hydrogen peroxide (30% H2O2 in H2O, CAS No. 7722-84-1),
hydrochloric acid (HCl, ~37%, CAS No. 7647-01-0), sodium bisulfate solution
(NaHSO3, CAS No. 7631-90-5) and hydroiodic acid (HI, containing no
stabilizer, CAS No. 10034-85-2, 57 wt.%), methyl methacrylate (MMA, CAS
No. 80-62-6, 99%) monomers, and 2,2′–Azobis(2–methylpropionitrile) (AIBN,
CAS No. 78-67-1, 98%) were purchased from Sigma–Aldrich Company Ltd.
L–ascorbic acid (LAA, CAS No. 50-81-7, ≥ 99%) was supplied by Alfa Aesar.
Ethanol (CAS No. 64-17-5, ≥ 95%) was obtained from Fluka.
Multi-walled carbon nanotubes (MWNTs) and graphene oxide (GO) were
purchased from Chengdu Organic Chemicals Co. Ltd (TIMESNANO) in China.
The MWNTs (Product Code: TNIM2) have a purity of over 90 wt. %, length of
30–50 μm, outer diameter of 8–15 nm, and specific area of 250–300 m2/g. The
GO powder (Product Code: TNGO) has a purity of over 99 wt. %, layers of
1–10, diameter of 0.5–3 μm, and thickness of 0.55–1.2 nm. It should be noted
that in this work, the GO was mostly synthesized from graphite powder by a
modified Hummers’ method (as described in Section 3.2.1), while the GO
purchased from Chengdu Organic Chemicals Co. Ltd (TIMESNANO) was only
used for a comparison of the morphology and multifunctional properties of the
57
as-prepared graphene aerogels (GAs).
All the chemicals were used as received without further purification.
3.2 Experimental techniques
3.2.1 Synthesis of graphene oxide (GO)
Graphene oxide (GO) was prepared from graphite powder by a modified
Hummers’ method [117, 253, 254]
. Typically, 2 g NaNO3 and 4 g graphite powder
were added to 100 ml concentrated H2SO4 in an ice bath. After stirring the
above mixture for 30 min, 14.6 g KMnO4 was then slowly added to the mixture
with stirring and cooling to keep the temperature lower than 20 °C for another 2
h. Next, the temperature of the mixture was increased to 35 °C with stirring for
12 h to convert graphite to graphite oxide. Following this, 180ml deionized (DI)
water was added to the mixture with stirring for 15 min, and 14 ml 30% H2O2
and 110 ml DI water were gradually added to the reaction mixture. Thereafter,
the color of the mixture changed from brown to yellow. The graphite oxide was
separated from the mixture by centrifugation and washed with 1 M HCl and DI
water to remove any impurities. The solid graphite oxide was finally dried at 60
°C for 3 days before collection.
To convert graphite oxide to graphene oxide (GO), the as-prepared
graphite oxide solid was first dissolved in DI water to form a 0.1–0.5 wt. %
aqueous suspension. The suspension was then sonicated (Sonics VCX-130
probe sonicator, 130 W, 100% amplitude) for 12 h to exfoliate the interlayers
of graphite oxide. GO powder was finally obtained by centrifuging the
sonicated suspension and drying the collected solid at 60° C for 3 days.
58
3.2.2 Synthesis of graphene aerogels (GAs)
GO aqueous suspensions with different concentrations of 1, 2, 6, and 12
mg/ml were prepared and placed in 20 ml vials with plastic caps. L–ascorbic
acid (LAA) was added to the GO suspension, with magnetic stirring for 5
minutes until it was completely dissolved. After that, the reaction mixture was
heated for a number of hours (5, 10, 20, 40 h) at different temperatures of 45,
70, and 95 °C. The as-prepared graphene hydrogels were immersed in DI water
for 3 days to remove excessive LAA before being placed into ethanol for
solvent exchange. The ethanol was changed every 24h for 3 days. Following
this, wet gels were subsequently dried with supercritical CO2 for duration of 45
min (3 cycles) to obtain the GAs.
To investigate the effects of reducing agents on the properties of the GAs,
a 2 mg/ml GO aqueous suspension was prepared. Three common reducing
agents (L–ascorbic acid, HI, and NaHSO3) were respectively added to the GO
suspension and sonicated until completely dissolved. The reaction mixture was
then heated at 95 °C for 5 h to form graphene hydrogels. The obtained graphene
hydrogels were immersed in DI water for 3 days to remove excessive reducing
agents, and then placed into ethanol for solvent exchange for another 3 days.
Finally, the wet gels were dried with supercritical CO2 to form the GAs.
In order to investigate the impact of thermal annealing on the electrical and
thermal properties of the GAs, the as-prepared GAs were also annealed at 450
°C for 5 h in an Argon (Ar) environment.
59
3.2.3 Synthesis of graphene–CNT hybrid aerogels
To develop graphene–carbon nanotube (CNT) hybrid aerogels, 10 mg
MWNTs were added to 10 ml DI water and dispersed with an ultrasonic
processor (Sonics, VCX 130). This solution was combined with 10 ml of 4
mg/ml GO suspension and sonicated for 4 h to obtain a homogenous GO–CNT
complex. Next, 160 mg LAA was added to the above suspension and sonicated
for 90 s to completely dissolve the LAA. The reaction mixture was heated at 95
°C for 5 h to form the graphene–CNT hybrid hydrogels. After solvent exchange
in DI water and ethanol, supercritical drying yielded the graphene–CNT hybrid
aerogels.
3.2.4 Synthesis of GA–poly (methyl methacrylate) (PMMA) nanocomposites
The GA–poly (methyl methacrylate) (PMMA) nanocomposites were
prepared via an in situ bulk polymerization process. Accordingly, 30 g methyl
methacrylate (MMA) and 20 mg 2,2′–Azobis (2-methylpropionitrile) (AIBN)
were first mixed in a 50 ml beaker and stirred at 75 °C for 30 min using a
magnetic stirrer. When the mixture became viscous, it was rapidly cooled
below 50 °C to cease the pre-polymerization. The GAs were then immersed into
the mixture and the mixture with the GAs immerged was kept at a low
temperature (below 10 °C) to remove air bubbles. After that, the above mixture
was maintained in a water bath at 50–55 °C for 30 h to cure the pre-polymerized
PMMA into a solid. The cured GA–PMMA nanocomposites were finally
machined and mechanically polished to remove the excessive part of PMMA
and achieve a smooth surface for subsequent characterization.
60
3.3 Characterization
3.3.1 X-ray diffraction (XRD)
X-ray diffraction (XRD) is one of the common techniques for examining
detailed information about the chemical composition and crystallographic
structure of a substance [255]
. With materials having a crystalline structure, their
ordered features will coherently scatter the X-rays in the directions that meet the
criteria for constructive interference, and thus lead to signal amplification. The
conditions required for constructive interference are determined by Bragg’s
law:
𝑛𝜆 = 2𝑑𝑠𝑖𝑛𝜃 (3.1)
where n is a integer, λ corresponds to the X-ray wavelength, d refers to the
distance between the lattice planes, and θ is the angle of incidence with the
lattice plane.
In this thesis, X-ray diffraction (XRD, 6000 Shimadzu, Japan) was
conducted to investigate the structures of graphite, GO and GA, and the effects
of GO concentrations and thermal annealing processes on the structures of GAs.
A diffractometer with a Cu-Kα radiation source (λ=0.1506 nm) was used. The
continuous scan was recorded from 5 ° to 30 ° (2θ) with a scan step of 0.02 ° and
a scan rate of 0.5 °/min.
3.3.2 Scanning electron microscopy (SEM)
Scanning electron microscopy (SEM) is a type of electron microscopy that
produces images displaying information about the surface morphology and
composition of a sample [256]
. The sample is scanned by a focused beam of
61
electrons, where the electrons interact with atoms in the sample and generate
various detective signals. Among them, the detection of secondary electrons
emitted by atoms is the most common mode. The secondary electrons are first
collected by inducing them towards an electrically biased grid, and then
accelerated towards a phosphor-biased one. Afterwards, the output signal of
these accelerated secondary electrons is amplified and shown as a 2D intensity
distribution, which can be eventually saved as a digital image after the
analog-to-digital conversion.
In this thesis, the nanostructures of GAs were characterized using field
emission scanning electron microscopy (FESEM, Model S-4300 Hitachi, Japan)
with a field-emission-gun operating at 10 kV for the GAs and 15 kV for the
GA–PMMA nanocomposites. Due to their good electrical conductivity, the GA
samples for FESEM were directly loaded on the holder without any pre-coating,
while the GA–PMMA nanocomposites were gold coated under 20 mA for 10s
to obtain better conductivity before surface scanning.
3.3.3 Physical adsorption/desorption of nitrogen (BET)
BET, which refers to the Brunauer–Emmett–Teller (BET) theory,
describes the physical adsorption of gas molecules on a solid surface, and is
considered the basic technique for analyzing the specific area of a material [257]
.
The concept of the BET theory originates from the Langmuir theory, but
extends adsorption of the monolayer molecular to that of the multilayer with
some specific hypotheses. These are: (i) gas molecules physically adsorb on a
solid in layers, infinitely; (ii) no interaction exists between each adsorption
62
layer, and (iii) each of these layers satisfies the Langmuir theory. The BET
equation is given below:
1
𝑣[(𝑝0 𝑝⁄ )−1]=
𝑐−1
𝑣𝑚𝑐(
𝑝
𝑝0) +
1
𝑣𝑚𝑐 (3.2)
where p and p0 are the equilibrium pressure and saturation pressure of
adsorbents at the test temperature, and v and vm are the volume of gas adsorbed
at equilibrium pressure and the monolayer capacity of the adsorbent
respectively. c is the BET constant and is expressed as:
c = exp (𝐸1−𝐸𝐿
𝑅𝑇) (3.3)
where E1 is the heat of adsorption for the first layer, while EL is that for the
second and upward layers, equal to the heat of liquefaction. R and T are the gas
constant and test temperature respectively.
Through the BET theory, the total surface area and specific surface area
are expressed as:
𝑆𝑡𝑜𝑡𝑎𝑙 =(𝑣𝑚𝑁𝑠)
𝑉 (3.4)
𝑆𝐵𝐸𝑇 =𝑆𝑡𝑜𝑡𝑎𝑙
𝑎 (3.5)
where vm is in the unit of volume, N is the Avogadro’s number, s is the
adsorption cross section, V the molar volume of the adsorbate gas, and a is the
net mass of the solid sample.
In this thesis, nitrogen adsorption/desorption measurements were carried
out with a Nova 2200e (Quantachrome) to obtain the pore properties of the
resulting GAs, such as BET specific surface area, BJH pore size distribution
and total pore volume. All the samples were degassed in a vacuum at 120 °C for
2 h before measurement.
63
3.3.4 Electrical conductivity measurement
Electrical conductivity measures the ability of a material to conduct
electrical current. The bulk electrical properties of the GAs or GA–PMMA
nanocomposites were measured using a two-probe method on a Solartron
1260+1287 electrochemical system, as shown in Figure 3.1 below. The samples
were sandwiched between two copper plate electrodes with silver paste applied
to eliminate the contact resistance. Except for the weight of the upper copper
electrode, no other weight was loaded on the top of the tested samples. The
dimensions of the samples were measured with a digital caliper before the
electrical conductivity measurement, which are 10–15 mm in diameter and 1–3
mm in thickness.
Figure 3.1 A Solartron 1260+1287 electrochemical system for electrical
property testing.
The electrical conductivities of the GAs and the GA–PMMA
nanocomposites were calculated from impedance values using the following
equation:
𝜎 =𝐼𝐿
𝑉𝑆 (3.6)
where V and I are the applied voltage and tested current through the samples,
64
and S and L are the cross section area and thickness of the tested samples,
respectively.
The electrical conductivity measurement was conducted within a voltage
range of 0–0.6 V, and each GA and GA–PMMA nanocomposite sample was
measured at least three times for consistency. The electrical conductivities of
the GAs and GA–PMMA nanocomposites were obtained by averaging the
values from the three tests.
3.3.5 Thermal conductivity measurement
The thermal conductivity of the GAs was measured using an improved
comparative infrared microscopy technique. Prior to the measurement, a GA
sample was sandwiched between two reference layers to build a 3-layered stack
using silver paste to ensure good contact between the sample and reference
layers, which was subsequently affixed to a heat sink plate at the bottom and a
resistive heater on the top. Amorphous quartz with a thermal conductivity of
kamorphous quartz = 1.3 W/m·K [258]
was selected as the reference material for this
thesis to ensure a comparable thermal resistance between the tested sample and
reference materials, as well as good thermal conduction through the entire
sample stack. The dimensions of the reference layers were 10 mm (L) × 10 mm
(W) × 1 mm (H), and the in-plane dimensions of the GA samples were also 10
mm (L) × 10 mm (W). All the surfaces of the GAs were polished and the
thickness of the GAs was controlled to be 1.2 ± 0.2 mm. A one-dimensional
heat flux was generated through the stack by the resistive heater and a
VariCAM high resolution thermographic system captured the temperature
65
distribution across the whole stack. All the surfaces facing the infrared (IR)
camera were coated with graphite to achieve a uniform and high (near unity)
emissivity. The temperature resolution of the IR camera was 0.08 K, and the
temperature through the whole stack was lower than 80 °C for all the
measurements.
The experimental set-up is illustrated in Figure 3.2(a) below. From the
measured two-dimensional temperature maps for the amorphous
quartz–GA–amorphous quartz stack, one-dimensional temperature profiles are
obtained by averaging the temperature in the direction perpendicular to the heat
flux. Figure 3.2(b) shows a typical temperature profile, where the temperature
gradients in the reference amorphous quartz region ((𝑑𝑇
𝑑𝑥)
amorphous quartz) and
sample region ((𝑑𝑇
𝑑𝑥)
GA) are calculated through fitting the temperature profile
with the least-square method [62]
.
Heat transport within the sample stack can be described by Fourier’s law.
Given the same cross section and constant heat flux in the stack, the
one-dimensional steady-state heat conduction equation is expressed as:
𝑞" = −kamorphous quartz ∙ (dT
dx)
amorphous quartz= −kGA ∙ (
dT
dx)
GA (3.7)
where 𝑞" is the heat flux through the sample stack and is calculated from the
value of kamorphous quartz and the average of the temperature gradients in the two
amorphous quartz regions. Accordingly, the thermal conductivity of the GA
(kGA) is determined from Equation (3.7) and the measured temperature gradient
in the GA.
In this thesis, to minimize the effect of the thermal boundary resistance
66
(TBR), the two pixels (~133 μm/pixel) at the boundary of each layer in the stack
were eliminated from the calculation [62]
. To yield robust results, the heat flux
(𝑞" ) versus the temperature gradient in the sample region ((𝑑𝑇
𝑑𝑥)
GA) was
plotted at several power levels, and the measured thermal conductivity of the
GAs (kGA) was extracted from the slope of the least-squares best fit to this curve,
as shown in Figure 3.2(c).
Figure 3.2 (a) Set-up schematic of thermal conductivity measurement using the
comparative infrared thermography technique. (b) Temperature distribution and
linear best fit curves for the 3-layered stack, which consists of a GA sample
sandwiched between two amorphous quartz layers, scale: 133 μm/ pixel. (c)
Heat flux as a function of temperature gradient in the GA region at several
power levels. During the measurement, the heat flux is typically below
11.25kW/m2.
67
3.3.6 Thermogravimetric analysis (TGA)
Thermogravimetric analysis (TGA) is a technique to determine thermal
characteristics, such as degradation temperatures, decomposition point, or the
contents of inorganic and organic components of a sample, according to its
weight changes as a function of temperature [259]
. The whole system consists of
a temperature-controlled furnace, a dual-range balance with two sample
containers, and attached thermal couples. During a typical test, approximately
10 mg powdered sample is located in one of the sample containers, while the
other acts as a reference. As the temperature of the furnace increases in a certain
gas environment, such as air, oxygen, nitrogen, or argon, the temperature of
each container and the mass of the sample are recorded accordingly.
Differential thermal analysis (DTA) is commonly performed simultaneously
with TGA, providing information about exothermic or endothermic reactions
that take place. In DTA, the temperature of the test sample is measured relative
to that of an adjacent inert material. It presents the difference in voltage
between the output of the sample thermocouple and the reference
thermocouple with temperature.
For this thesis, the TGA and DTA of both GAs and GA–PMMA
nanocomposites were conducted on a DTG60H thermogravimetric analyzer
from room temperature to 1000 °C in air with a heating rate of 5 °C/min.
3.3.7 Vickers microhardness measurement
The hardness of a material is defined as its resistance to indentation, and it
is usually determined by measuring the permanent depth of the indentation.
68
Microhardness testing, also refers to as microindentation, is a common method
for measuring the hardness of a material, especially for small parts and thin
sections, on a microscopic scale [260]
.
The Vickers hardness measurement is based on an optical measurement
system and follows the procedures of ASTM E-384, which is commonly used to
determine and verify the Knoop and Vickers hardness of materials. During the
measurement, a square-based pyramid indenter is used to make an indentation
on the testing specimen under a light load. The hardness value is calculated
according to the length of the indentation, which could be precisely measured
by the optical microscopy of the testing machine. The calculation is given in the
following equation:
𝐻𝑉 =𝐹
𝐴≈
0.1891𝐹
𝑑2 (3.8)
where F is the applied force, [N], A is the surface area of the indentation, [mm2],
and d is the mean diagonal length of the indentation, [mm]. The calculated
hardness HV is in GPa.
For this thesis, the Vickers microhardness was measured with a
Shimadzu-HMV automatic digital microhardness tester using an individual
indentation force of 245.2 mN for duration of 15 seconds.
69
CHAPTER 4: Morphology Control of Graphene
Aerogels
4.1 Introduction
In recent years, a large amount of research has been conducted to develop
graphene or graphene-based materials for various applications, such as energy
storage devices, sensors, and nanocomposites. It is believed that the realization
of a three-dimensional graphene nanostructure would represent a significant
step towards fulfilling the potential of graphene. However, as presented in
Section 2.2, the methods reported thus far for the fabrication of 3D graphene
nanostructures are still very limited, and a binder is usually required to achieve
the assembly. Although the chemical reduction method takes advantages of
binderless and more environmentally friendly than other approaches, very few
studies have been conducted to investigate the effects of its synthesis
parameters on the morphology and multifunctional properties of the
as-prepared GAs. The effect of different reducing agents on the GA
nanostructures has also not been reported.
In addition, a one-of-a-kind symbiotic relationship has been observed to
exist between one-dimensional CNTs and two-dimensional graphene sheets.
Therefore, integrating CNTs into the GAs could further enhance the electrical
and mechanical properties of the GAs [157, 212]
. Nevertheless, the effect of CNTs
on the morphology of GAs has also yet to be studied.
In this thesis, we report systematic research to effectively control the
morphology of GAs. The GAs were synthesized through a simple chemical
70
reduction method. Since no polymeric binder was applied during the process,
the developed method is effective and environmentally friendly for mass
production of GAs. As the nanostructures of the GAs can be easily controlled
by adjusting the synthesis parameters, such as compositions (GO
concentrations), reduction temperatures, reduction times, reducing agents, and
CNT inclusions, these experimental results could serve as a guide to synthesize
the GAs with the desired nanostructures, which would be very useful for the
electrodes of energy storage devices such as supercapacitors and lithium
batteries, and nanocomposites.
4.2 Nanostructured control of graphene aerogels (GAs)
GAs are synthesized through a modified Hummers’ method and chemical
reduction method as described in Section 3.2.2. The mechanism of graphene
self-assembly by chemical reduction was first proposed by Chen et al. [145]
It is
widely known that GO can be well dispersed in DI water due to its
hydrophilicity. When the GO aqueous suspension is reduced by reducing agents
such as LAA, NaHSO3 and HI, the reduced GO (rGO) nanosheets become
hydrophobic and the π–π interaction between these hydrophobic rGO
nanosheets significantly increases, which finally leads to the formation of a 3D
structure of graphene.
Except when investigating the effects of reducing agents on the
multifunctional property of GAs, LAA is selected as the major reducing agent
for the other sections of this thesis, because it is efficient and environmentally
friendly for the reduction of GO to graphene; the excess LAA can also be easily
71
removed during the solvent exchange by DI water. According to the estimated
balanced chemical reaction reported elsewhere, the required amount of LAA
should be at least 3.3 times that of the GO content by mass [157, 261]
. Thus the
amount of LAA used for this work is 4 times that of the GO content by mass, to
ensure a sufficient reduction. Figure 4.1 shows (a) the aqueous suspension of
GO and a graphene hydrogel and (b) a graphene aerogel.
Figure 4.1 (a) Aqueous suspension of GO and a graphene hydrogel and (b) A
graphene aerogel.
The oxidation of graphite into GO (or GO after exfoliation) and the
reduction of a GO into a GA are investigated by X-ray diffraction (XRD)
technique, as shown in Figure 4.2 below. From the XRD spectrum, a strong
characteristic peak of 26.4° (JCPDS card number: 75-1621) is observed for the
raw graphite powder, while a peak of 11° is found for the graphene oxide
obtained from the oxidation and exfoliation of graphite, which is consistent
with the results reported elsewhere and indicates the conversion of graphite
(Interlayer spacing: 0.34 nm) to graphite oxide (Interlayer spacing: 0.37 nm)
[14, 156, 203, 262]. For the XRD pattern of GA, a broad peak is shown at 23.5°,
indicating the effective removal of oxygen groups during the reduction process
and an amorphous structure formed between the graphene layers [145, 174]
.
Previous Raman spectra of the chemically derived graphene nanosheets
72
usually exhibit a D to G band intensity ratio (ID/IG) of ~0.99–1.43 (1.07 in this
work), indicating the introduction of defects after the oxidation and reduction
of graphite [263, 264]
.
Figure 4.2 XRD patterns of graphite, graphite oxide and graphene aerogel.
4.2.1 Effects of chemical compositions and synthesis conditions
The effects of chemical compositions and synthesis conditions, such as
reduction temperatures and reduction times, on the morphologies of the GAs
are systematically investigated, as shown in Table 4-1. The GAs are synthesized
with various initial GO concentrations of 1, 2, 6, and 12 mg/ml, reduction
temperatures of 45, 70, and 95 °C, and reduction times of 5, 10, 20, and 40 h,
respectively. The nanostructures of the GAs obtained under various GO
concentrations, different reduction temperatures, and different treatment times
are compared in Figure 4.3 below. All of them present interconnected
three-dimensional network structures. It can be seen that the 3D network
formed from 12 mg/ml GO suspension is denser than that formed from the
concentration of 2 mg/ml, while the GA synthesized at 95 °C has a more
packed structure than those at 70 and 45 °C. The GA sample reduced for 40 h
73
is also more overlapped than that for 5 h. This trend is also associated with
that of the surface areas of the GAs.
Table 4-1 Nanostructure control of the GAs.
No. Synthesis
conditions
Surface
area
(m2/g)
Average
pore size
(nm)
Pore
volume,
(cm3/g)
Density,
(mg/cm3)
Porosity*
(%)
GO concentrations (Temperature: 70 °C, time: 40 h)
1 1 mg/ml 284 3.6 0.51 18.24 99.13
2 2 mg/ml 302 4.1 0.63 26.53 98.74
3 6 mg/ml 365 3.9 0.72 56.44 97.31
4 12 mg/ml 473 4.0 0.95 67.97 96.76
Reduction times (GO concentration: 2 mg/ml, temperature: 70 °C)
5 5 h 560 3.9 1.09 18.47 99.12
6 10 h 438 4.0 0.88 20.81 99.01
7 20 h 381 3.9 0.74 22.84 98.91
2 40 h 302 4.1 0.63 26.53 98.74
Reduction temperatures (GO concentration: 2 mg/ml, time: 40 h)
8 45 °C 577 4.0 1.17 25.61 98.78
2 70 °C 302 4.1 0.63 26.53 98.74
9 95 °C 245 3.8 0.46 35.78 98.30
*Porosity is converted from density ratio of the GAs to graphene (2.10 g/cm3).
Figure 4.3 SEM images of graphene aerogels synthesized under different
conditions: (a) 2 mg GO/ml, 70 °C, 40 h; (b) 12 mg GO/ml, 70 °C, 40 h; (c) 2
mg GO/ml, 45 °C, 40 h; (d) 2 mg GO/ml, 95 °C, 40 h; (e) 2 mg GO/ml, 70 °C, 5
h. Compared with image (a), more graphene nanosheets can be observed in
image (b). The GA sample in image (c) shows a more loose structure than those
in image (a) and image (d). Image (e) also shows a less dense network than that
in image (a).
74
The characterization of the specific surface areas and porous properties of
the GAs obtained under different conditions is conducted by nitrogen
adsorption/desorption measurements with a Nova 2200e (Quantachrome).
Before the measurements, samples are degassed at 120 °C in a vacuum for 2 h to
remove moisture, and the calculations of the specific surface area and pore size
of the GA samples are conducted by the Brunauer–Emmett–Teller (BET) and
Barrett–Joyner–Halenda (BJH) methods. A typical nitrogen
adsorption/desorption isotherm is shown in Figure 4.4. Type IV
adsorption/desorption isotherms are displayed for all the GA samples,
indicating that there are many mesopores in the GA structures. All the
morphological characteristics of the as-prepared GAs are summarized in Table
4-1 above.
Figure 4.4 A typical N2 adsorption/desorption isotherm of a GA.
For all the GAs synthesized at 70 °C for 40 h, the highest surface area of
473 m2/g is obtained with 12 mg/ml initial GO concentration, while the lowest
value of 284 m2/g is found with a GO concentration of 1 mg/ml. It can be seen
75
that under the same reduction conditions, the surface area and pore volume of
the GAs increase along with the increase in the initial GO concentration, since
more graphene nanosheets are more likely to create more pores during the
reduction process. The effect of synthesis conditions on GAs is tested with 2
mg/ml GO throughout. With the same reduction time of 40 h, the GA
synthesized at 45 °C has the largest surface area of 577 m2/g and the largest pore
volume of 1.17 cm3/g, while that at 95 °C has the lowest values of 245 m
2/g and
0.46 cm3/g, respectively. When the reduction temperature is kept at 70 °C, the
surface area and pore volume values of the GA sample with 5 h reduction
treatment are much larger than those of the one treated for 40 h.
During the reduction, the formation of hydrophobic sites and π–π
interactions between graphene sheets could be the main factor determining the
porous structure of the GAs [174]
. In this case, a higher temperature or a longer
time may probably help the reduction process to be more efficient or more
complete, thus leading to the creation of more overlapping sites. Under a higher
reduction temperature, the higher pressure in the reaction vial may also favor
the assembly of graphene nanosheets. Such assembly and overlapping between
the graphene nanosheets consequently results in a reduction of the surface area
and pore volume of the as-prepared GAs. Therefore, it can be concluded that
with the same GO concentration, the surface area of the GAs is reduced with the
increase in reduction temperature or reduction time. Compared with the
previous work [174]
, the surface area and pore volume of the GAs prepared by
the mild chemical reduction method are all higher than those of the samples
prepared by the hydrothermal method, indicating the negative effect of higher
76
pressure on the surface area and pore volume of the GAs.
4.2.2 Effects of reducing agents
Except for synthesis conditions, the selection of reducing agent (LAA, HI,
or NaHSO3) also impacts the morphology, electrical property, and thermal
stability of the GAs. The effect of reducing agents on the morphology of GAs is
investigated when 2 mg/ml GO suspension is reduced by LAA, HI, and
NaHSO3 under the same synthesis conditions; the amounts of three reducing
agents are enough for a sufficient reduction [145, 157]
. The GA morphologies are
characterized by nitrogen adsorption/desorption measurements to show the
specific surface area and pore size, obtained by the Brunauer–Emmett–Teller
(BET) and Barrett–Joyner–Halenda (BJH) methods, respectively. Table 4-2
summarizes the morphologies of the GAs reduced by LAA, HI, and NaHSO3,
with and without thermal annealing conditions.
Table 4-2 Effects of reducing agent and addition of CNTs on morphology of
the GAs.
Thermal
annealing
Surface area
(m2/g)
Pore size
(nm)
Density
(mg/cm3)
Porosity
(%)
Reducing agents
LAA Without 423 3.6 28.1 98.66
With 487 3.6 28.2 98.66
HI Without 385 3.6 18.4 99.12
With 281 3.6 16.4 99.22
NaHSO3 Without 679 3.6 20.2 99.04
With 582 3.6 22.3 98.94
CNTs concentration (mg/ml) (LAA as the reducing agent)
0 – 738 3.6 37.0 98.24
0.5 – 844 3.6 31.2 98.51
1 – 788 3.6 36.1 98.28
2 – 636 3.6 41.3 98.03
The surface areas of the GAs reduced by LAA, HI, and NaHSO3 are 423,
77
385 and 679 m2/g respectively, due to the different reduction effectiveness of
the three reducing agents. The more efficient reducing agent (e.g. HI) removes
more functional groups (–OH, –COOH, –O–, etc.) during the reduction
processes and forms more π–π bonding and stacks between graphene
nanosheets, which decreases the surface area of the GAs.
After thermal annealing of the GAs, the surface areas of the GAs reduced
by HI and NaHSO3 decrease to 281 and 582 m2/g, respectively. The annealing
process further removes the functional groups of the GAs, causing them to
condense further. More π–π stack formation between the graphene sheets
causes decreased GA surface area. However, the surface area of the GAs
reduced by LAA increased to 487 m2/g after thermal annealing. There is no
clear explanation to explain the increase in surface area for the LAA–reduced
GAs, although it is possible that the thermal annealing process may decrease the
π–π interaction among graphene sheets and dehydroascorbic acid (the oxidized
form of LAA), in addition to removing the functional groups on graphene
sheets.
4.2.3 Effects of CNTs
As explained previously, graphene hydrogels are formed by π–π stacking
during the reduction of hydrophilic GO to hydrophobic graphene nanosheets.
The mechanism of graphene–CNT hybrid aerogel formation may be also due to
the π–π interaction between graphene nanosheets, as well as that between
graphene and CNTs. During the reduction of GO in its aqueous suspension, the
majority of the CNTs are physically wrapped inside the 3D network of
78
graphene hydrogel, while a minority of the CNTs may form a π–π interaction
with the reduced graphene oxide nanosheets. Figure 4.5(a) shows an
as-prepared graphene–CNT hybrid aerogel using LAA, with 10 mm diameter
and 3 mm thickness.
Figure 4.5 (a) A digital photo of a graphene–CNT hybrid aerogel; (b) and (c)
SEM images of a graphene–CNT hybrid aerogel at different magnifications.
Different concentrations of CNTs (0, 0.5, 1 and 2 mg/ml) are added into
the GAs to form graphene–CNT hybrid aerogels. Figures 4.5(b) and (c) show
SEM images of the interconnected 3D network within the graphene–CNT
hybrid aerogel, where the CNTs are dispersed inside the graphene sheet
network. As seen in Table 4-2 previously, the surface area of the
graphene–CNT hybrid aerogels increases with CNT concentrations up to 1
mg/ml. The surface area of the graphene–CNT hybrid aerogels increases
because of: (i) the additional surface area of the dispersed CNTs, and (ii)
separation of the graphene nanosheets by the CNTs during the self-assembly
79
processes. However, when the CNT concentration is further increased to 2
mg/ml, the surface area of the graphene–CNT hybrid aerogel decreases to less
than that of the pure GAs. At high CNT concentrations, local agglomeration and
bundling of the CNTs may block the pores of the GAs [265]
.
4.3 Conclusions
In summary, GAs are synthesized by a simple chemical reduction method
with various chemical compositions (initial GO concentrations), different
synthesis conditions (reduction temperatures and reduction times), and three
reducing agents (LAA, HI, and NaHSO3). A morphology control of the
as-prepared GAs can be achieved by changing these synthesis parameters
appropriately. The results show that the GAs synthesized from a higher GO
concentration possess higher surface areas, while higher reduction temperature
and reduction time make GAs a packed structure with lower surface areas.
Similarly, the GAs synthesized with a more efficient reducing agent exhibit a
lower surface area, while those synthesized with a low-efficiency reducing
agent possess a higher surface area. In addition, graphene–CNT hybrid aerogels
were successfully developed by adding CNTs during the GA synthesis. The
addition of CNTs enhances the surface area by up to 14%, compared with the
pure GAs. These results provide insight into optimizing the nanostructures of
the GAs for various applications, including the development of electrodes,
energy storage devices and nanocomposites.
The nanostructure control of GAs will be summarized again in the
Conclusion chapter.
80
CHAPTER 5: Multifunctional Properties of
Graphene Aerogels
5.1 Introduction
Chapter 4 presents systematic research on the effective control of the
morphology of GAs. However, this morphological effect also has a significant
impact on the multifunctional properties of the GAs. In this chapter, the
electrical property, thermal property, and thermal stability of the GAs are
comprehensively investigated in relation to their variation in morphologies.
These experimental results are thus useful for guiding the optimization of the
surface area and multifunctional properties of the GAs at the same time.
For this thesis, the electrical property and thermal stability of the GAs were
measured by a two-probe method and TGA, respectively. In terms of the
measurement of thermal property, a comparative infrared (IR) microscopy
technique was developed based on the ASTM E1225 standard. Due to its
satisfactory accuracy and facile operation, this IR microscopy technique was
applied to determine the thermal conductivities of the GAs fabricated from
various graphene oxide (GO) aqueous suspensions. The thermal conductivity
was measured as a function of GO concentration and thermal annealing. This
thesis reports the first benchmark data of the thermal properties of the GAs, and
seeks to identify the factors limiting their thermal conductivities.
81
5.2 Electrical property of graphene aerogels (GAs)
5.2.1 Effects of chemical compositions and synthesis conditions
Figure 5.1 below shows a typical current–voltage curve of the GA, which
is obtained from a 12 mg/ml GO suspension at 70 °C for 40 h. Its electrical
conductivity is calculated from this linear I–V plot based on the Ohm’s law. It
can be observed from Table 5-1 below that under the same reduction
temperature and reduction time, GA synthesized from a higher initial GO
concentration has a higher electrical conductivity, i.e. the highest electrical
conductivity of 2.1 S/m is obtained when the initial GO concentration is 12
mg/ml, which is comparable to the data reported elsewhere [266]
. A likely
explanation is that more graphene nanosheets in a higher GO concentration can
form more overlapping during the synthesis process, thus attaining better
electron mobility.
Figure 5.1 I–V plot of the GA obtained from 12 mg/ml GO suspension at 70 °C
for 40 h.
82
Table 5-1 Morphological effects on electrical property of the GAs.
No. Synthesis
conditions
Surface area
(m2/g)
Pore volume,
(cm3/g)
Conductivity,
(S/m)
GO concentrations (Temperature: 70 °C, time: 40 h)
1 1 mg/ml 284 0.51 0.7
2 2 mg/ml 302 0.63 0.8
3 6 mg/ml 365 0.72 1.1
4 12 mg/ml 473 0.95 2.1
Reduction times (GO concentration: 2 mg/ml, temperature: 70 °C)
5 5 h 560 1.09 0.5
6 10 h 438 0.88 0.6
7 20 h 381 0.74 0.7
2 40 h 302 0.63 0.8
Reduction temperatures (GO concentration: 2 mg/ml, time: 40 h)
8 45 °C 577 1.17 0.5
2 70 °C 302 0.63 0.8
9 95 °C 245 0.46 1.0
Figure 5.2 shows the effect of initial GO concentrations on the electrical
conductivities and surface areas of the GAs. It may be noted that increasing the
concentration of the initial GO aqueous suspension can increase both the
electrical conductivity and surface area of the GA. However, if the GO
concentration is too high, it will be very difficult to form a uniform structure in
the as-prepared GA. For GAs synthesized from the same initial concentration of
GO, the electrical conductivities decrease with the increase in surface areas.
This is mainly because for a denser structure of GA, there would be more
stacking of graphene nanosheets and thus stronger π–π interaction existing in
its network. It has been reported that the electrical conduction of chemically
reduced graphene oxide nanosheets is dominated by an electron hopping
mechanism. Therefore, as more graphene nanosheets become intact, the
electron hopping is expected to be significantly enhanced, thus leading to a
reduction of the contact resistance and a higher electrical conductivity.
83
Figure 5.2 Effects of initial GO concentrations on electrical conductivity and
surface area of GAs.
5.2.2 Effects of reducing agents
The HI–reduced GAs have the highest electrical conductivity of 3.29 S/m,
while the LAA and NaHSO3–reduced GAs have relatively lower electrical
conductivities of 2.08 and 1.67 S/m, respectively (see Table 5-2 and Figure 5.3
below). The differences in the electrical conductivities of the GAs reduced by
different reducing agents can also be explained by the morphological
differences. The GA with a lower surface area has a higher electrical
conductivity due to the more densely packed nanostructure, which enhances the
electron transport.
Table 5-2 Effects of reducing agent and addition of CNTs on electrical
conductivity and thermal stability of the GAs.
Thermal
annealing
Surface
area (m2/g)
Electrical
conductivity
(S/m)
DTA peak
temperature (°C)
Reducing agents
LAA Without 423 2.08 ± 0.20 539
With 487 4.12 ± 0.53 568
HI Without 385 3.29 ± 0.44 591
With 281 5.90 ± 0.04 598
NaHSO3 Without 679 1.67 ± 0.03 540
With 582 9.52 ± 0.38 542
84
CNT concentration (mg/ml) (LAA as reducing agent)
0 – 738 0.20 ± 0.01 516
0.5 – 844 0.33 ± 0.03 495
1 – 788 0.56 ± 0.09 496
2 – 636 0.47 ± 0.06 532
Figure 5.3 Effect of reducing agents and thermal annealing on (a) electrical
conductivity and (b) thermal stability of the GAs.
85
Figure 5.4 below shows the effects of the thermal annealing on the
structures and properties of the GAs. The XRD spectrum for all the GAs shows
a broad peak around 23.5–25 ° before annealing, which is consistent with
previous reports [14, 178, 203, 262]
and indicates the formation of non-crystalline
structures by the π–π stacking among the graphene nanosheets. For the annealed
GAs, the XRD spectrum shifts to the right, indicating the formation of a more
condensed nanostructure after the thermal annealing treatment. During thermal
annealing, more functional groups may be removed and stronger π–π
interactions could form among the graphene sheets. Therefore, the annealed
GAs have much higher electrical conductivities than those of the non-annealed
GAs, as shown in Figure 5.3(a) above.
Figure 5.4 XRD patterns of the GAs reduced by different reducing agents with
and without thermal annealing.
5.2.3 Effects of CNTs
The addition of CNTs into the GAs also enhances their electrical
conductivities, as shown in Table 5-2 previously and Figure 5.5 below. When
86
0.5 mg/ml CNTs is added into the GA, the electrical conductivity of the
graphene–CNT hybrid aerogel is 0.33 S/m, a 65% increase over that of the pure
GA. When the concentration of CNTs is 1 mg/ml, the electrical conductivity of
the graphene–CNT hybrid aerogel reaches a maximum of 0.56 S/m, while 2
mg/ml CNTs slightly decrease the electrical conductivity of the graphene–CNT
hybrid aerogel to 0.47 S/m. Within the graphene–CNT hybrid aerogels, CNTs
could create more channels for efficient electron transfer between graphene
sheets and therefore enhance their electrical conductivities. However, when the
CNT concentration exceeds 1 mg/ml, the local agglomeration of CNTs reduces
the electrical conductivities of the GAs.
Figure 5.5 Electrical conductivity and DTA peak temperature for
graphene–CNT hybrid aerogels with various CNT concentrations.
5.3 Thermal stability of graphene aerogels (GAs)
5.3.1 Effects of chemical compositions and synthesis conditions
The thermal durability, which is an important factor for electrode materials
working at high temperatures, is studied by TGA on DTG60H using the GAs
synthesized at 45 and 95 °C for 40 h, respectively. Their TGA profiles are
87
displayed in Figure 5.6 below; both exhibit similar trends, but with different
burning temperatures. The GA obtained at 95 °C shows an obvious weight loss
in the range of 450–560 °C with a DTA peak at 544 °C, while the GA obtained
at 45 °C possesses a continuous decrease in weight from 200 to 570 °C with a
DTA peak at 529 °C. It can be seen that the GA synthesized at a higher
temperature creates a denser structure and higher thermal stability, which also
indicates that more overlapping of graphene nanosheets makes the GA a more
rigid network. For comparison, the GA synthesized at a lower temperature has a
looser structure and appears to be less thermally stable due to the existence of
some unpacked part inside.
Figure 5.6 Thermal stability of GAs at (a) 45 °C and (b) 95 °C.
5.3.2 Effects of reducing agent and CNTs
As shown in Figure 5.3(b) previously, the HI–reduced GA shows a DTA
peak at 598 °C, which is higher than that of the LAA– and NaHSO3–reduced
GAs (approximately 540 °C). During TGA tests, the remaining functional
groups on the graphene nanosheets are first removed before the GA reacts with
the oxygen in an air atmosphere. Consistent with the morphological
88
characteristic as discussed previously, the increased thermal stability of the
HI–reduced GA is likely due to its more condensed nanostructure and lower
surface area [145]
.
As shown in Figure 5.5 and Table 5-2 previously, the thermal stability of
GAs is affected by the CNT concentrations and the morphology of the
graphene–CNT hybrid aerogels. When the dispersed CNT concentration is
increased up to 1 mg/ml, the surface areas of GAs increase and their thermal
stabilities decrease due to the more porous structures. When the CNTs form
agglomeration/bundles, the surface area of the GAs is reduced and the thermal
stabilities are larger due to the more condensed nanostructures and the stronger
bonding between CNTs–CNTs and graphene–graphene nanosheets.
5.4 Thermal conductivity of graphene aerogels (GAs)
5.4.1 Effects of chemical compositions
During the reduction of hydrophilic GO to hydrophobic graphene, the
partial overlap of flexible graphene nanosheets due to π–π interaction results in
the formation of the 3D graphene hydrogels [14, 157]
. After supercritical drying,
the interconnected porous network is preserved in the GAs, as shown by the
FESEM images in Figure 5.7. The concentration of the GO aqueous suspension,
as well as the annealing treatment, impacts the morphologies and thermal
conductivities of the GAs. Table 5-3 summarizes the synthesis conditions,
density, volume fraction and thermal conductivities for each GA.
89
Figure 5.7 SEM images of the GAs synthesized under different conditions:
(a)–(h) represent GA1–GA8 respectively. All images present interconnected
3D networkstructures. With annealing, images (e)–(h) show denser networks
than those in images (a)–(d) at the same GO concentration.
90
Table 5-3 Synthesis conditions, density, volume fraction and thermal
conductivity of the GAs.
No.
GO
conc.
(mg/ml)
Thermal
annealing
Density
(mg/cm3)
Volume fraction
(vol. %)*
Thermal
conductivity kGA
(W/m·K)
GA1 1
Without
14.1 0.67 0.12 ± 0.006
GA2 2 28.1 1.34 0.18 ± 0.013
GA3 6 42.8 2.04 0.24 ± 0.007
GA4 12 52.4 2.50 0.36 ± 0.015
GA5 1
With
16.4 0.78 0.18 ± 0.006
GA6 2 28.2 1.34 0.23 ± 0.023
GA7 6 40.0 1.90 0.31 ± 0.023
GA8 12 49.0 2.33 0.28 ± 0.016 *Volume fraction is converted from density ratio of the GAs to graphene
(ρGraphene=2.10 g/cm3) using the equation: Vf= ρGA/ρGraphene×100%
[183, 208]. The
density of GA (ρGA) is determined by the mass and volume of each GA
sample.
The total thermal conductivity of the GAs consists of both electrical and
lattice contributions [183]
. However, the electronic contribution part kGA,e is
estimated to be less than 3% of the total thermal conductivity across the whole
measured temperature range, according to the Wiedemann-Franz law [267]
.
𝑘GA,e = 𝐿𝑇 𝜌GA⁄ (5.1)
where L, T, and ρ are the Lorentz number, temperature and electrical resistivity
respectively. Therefore, the lattice vibrations (phonons) are proposed to be the
dominant heat transfer mechanism in the GAs [55, 183, 194, 248, 268]
.
Our previous investigations [173, 174]
show that the GA synthesized from a
higher GO concentration exhibits higher electrical conductivity, due to the
better electron mobility through the contacts between graphene nanosheets.
Similar trends are also observed for the thermal properties of the GAs, as shown
in Figure 5.8 below. The thermal conductivity of the GA with 2.5 vol. %
graphene is 0.36 ± 0.015 W/m·K, which is significantly higher than those for
GAs with 0.67–2.04 vol. % graphene (0.12 ± 0.006, 0.18 ± 0.013 and 0.24 ±
91
0.007 W/m·K). This significant increase in thermal conductivity with the initial
GO concentration could be attributed to a more connected network of graphene
sheets in the GAs synthesized from a higher GO concentration (see Figure 5.7
(a)–(d)). Specifically, the higher density of graphene nanosheets and the
increased overlap between nanosheets is expected to provide lower resistivity
pathways for phonons to transport through the GAs [248]
.
Figure 5.8 Effects of GO concentrations and thermal annealing on thermal
conductivities of the GAs.
Considering the extremely high thermal conductivity of graphene (up to
5300 W/m·K), all the measured values of the GAs are extremely low. The GAs
have extremely high porosities (≥ 97.5%), and all the pores inside are filled by
air, which has very low thermal conductivity (0.026 W/m·K [269]
). The low
density of graphene is expected to be the dominant factor which results in the
low effective thermal conductivity of the GAs in the bulk form [270]
. However,
the measured thermal conductivity was still lower than the values predicted by
the porosity and the conductivity of the individual graphene nanosheets. Thus
additional factors must contribute to and reduce the thermal conductivities of
92
the GAs.
First, the quality of the graphene nanosheets significantly impacts the
thermal conductivity of the GAs [271]
. All the GAs studied for this thesis are
synthesized by the chemical reduction method, in which a number of defects are
introduced during the strong oxidation process and are not completely repaired
during the chemical or thermal reduction processes [46, 144, 272]
. Note that the
highest thermal conductivity of reduced graphene oxide (rGO) has been
measured to be only 6.8 ± 0.08 W/m·K [144]
, which falls well below the 5300
W/m·K of the defect-free graphene.
Second, the size of chemically derived graphene nanosheets is comparable
to the phonon mean free path (~775 nm at room temperature) [247]
. Given that
the individual nanosheets are reported to range in size from 0.2 to 2 μm [144, 273]
,
a large fraction of the graphene nanosheets within the GAs would be far smaller
than the mean free path (~775 nm). Therefore, their thermal conductivities
should also be significantly limited by the increased scattering at the nanosheet
boundaries [273]
.
Third, the GAs consist of numerous randomly distributed graphene
nanosheets in contact with each other. Heat must transfer between several
nanosheets as it conducts across the aerogel, and the thermal resistance between
the individual graphene nanosheets is likely quite high. This interface resistance
may limit the thermal conduction through the aerogel and increase its effective
thermal resistance. However, since the graphene nanosheets bridge the whole
GA bulk, the 3D networks still achieve more effective heat transport than
dispersed rGO in composites [183, 192]
.
93
Furthermore, the measurement technique may involve errors arising from
the limited resolution of the IR camera, the convection due to high surface areas
of the GAs, and variations in the thermal conductivities of the GAs and
reference materials with temperature. All these errors should be taken into
consideration and minimized during the experiments.
5.4.2 Effects of thermal annealing
As shown in Table 5-3, after thermal annealing, the thermal conductivities
of GA5–7 are enhanced (kGA=0.18 ± 0.006, 0.23 ± 0.023 and 0.31 ± 0.023
W/m·K, respectively), compared with GA1–3 without annealing. The thermal
conductivity of GA8 (0.28 ± 0.016 W/m·K) is slightly lower than that of the
sample without annealing (GA4), and is also lower than that of GA7, which is
annealed but has a lower GO concentration.
Post growth annealing treatments of carbon nanotubes (CNTs) or graphene
remove the residual functional groups and repair some defects [145, 146]
. The
improved thermal conductivities of GA5–7 can be attributed to the significant
elimination of saturated sp3 bonds bearing functional groups (which cause
enhanced phonon scattering and hinder the thermal transport [144, 271]
) and the
formation of sp2-hybridized carbon atoms. In contrast, the decreased thermal
conductivity of the GA8 may be due to the condensation of graphene
nanosheets during thermal annealing, as is indicated by the right-shift of the
XRD spectrum for GA8 in Figure 5.9. Phonon scattering between multi-layered
graphene sheets leads to a significant decline of the instinct thermal
conductivity.
94
Figure 5.9 XRD patterns of the GA4–5, 7 and 8. Broad peaks indicate the
non-crystal aerogel stucture. Compared with GA7, the peak of GA8 shifts right
and indicates the smaller interspacing between graphene nanosheets.
5.4.3 Prediction of thermal boundary resistance and thermal conductivity of
graphene nanosheets within GAs
The rule of mixtures [274]
and effective medium theory (EMT) [251]
have
been applied to predict the range of the thermal conductivity of the rGO
nanosheets, as shown in Figure 5.10 below. The GAs are considered a
two-phase system consisting of rGO and air with thermal conductivities kG and
kair, respectively (kair =0.026 W/m·K [269]
). A first estimation of the thermal
conductivity of graphene aerogels (kGA) can be calculated through the rule of
mixtures:
𝑘𝐺𝐴 = 𝑓 𝑘𝐺 + (1 − 𝑓)𝑘𝑎𝑖𝑟 (5.2)
where f is the volume fraction of graphene. This model predicts the thermal
conductivity of the rGO kG to be 12.2 W/m·K, neglecting the effects of the
thermal boundary resistance between nanosheets. Both the rule of mixtures and
the EMT approximations do not consider the details of the microstructure of the
95
GAs (pore size, shape etc.) or convective heat transfer within pores, but both
provide an estimation of the thermal conductivity of the graphene fraction of the
aerogel.
Figure 5.10 The thermal condictivity of the GAs as a function of graphene
volume fraction. Best fits from rule of mixtures (dashed line) and EMT (orange
dots) are shown. For the rule of mixtures, the slope of the least-squares best fits
to the experimental data is the (kG-kair). For EMT, the values calculated assume
a kG=30.2 W/m·K, a value which is obtained by manually fitting the data.
A modified effective medium theory (EMT) based model can also be used
to estimate the thermal conductivity:
𝑘𝐺𝐴 = 𝑘𝐺 [3𝑘𝑎𝑖𝑟+2𝑓(𝑘𝐺−𝑘𝑎𝑖𝑟)
(3−𝑓)𝑘𝐺+𝑘𝑎𝑖𝑟𝑓+𝑅𝐵𝑘𝑎𝑖𝑟𝑘𝐺𝑓
𝐻
] (5.3)
where RTBR is the TBR [250]
between rGO and air (RTBR=10-5
Km2/ W
[147, 275])
and H is the total thickness of rGO, taken to be 3 nm in this calculation. The
thickness of an rGO monolayer has been reported to be 0.6 to 0.9 nm [172, 276]
,
and here the GAs are assumed to consist of 5-layer rGO nanosheets on average.
This model assumes that the graphene nanosheets are randomly dispersed and
not interconnected. It should be noted that although RTBR is incorporated in the
96
EMT model, it does not take into account the TBR between nanosheets, but
rather the resistance between the nanosheets and air. The thermal conductivity
of the rGO of kG=30.2 W/m·K is determined from this EMT model [194, 277]
and
is considered an upper bound to the thermal conductivity. Thus the thermal
conductivity of the chemically derived graphene nanosheets in our work is in
the range of 12.2–30.2 W/m·K, which is a factor of 1.8–4.4 higher than the
value of the previous work [144]
.
5.5 Conclusions
In conclusion, the effects of chemical composition, synthesis conditions,
reducing agents and CNT inclusions on the multifunctional properties of the
GAs were systematically investigated in terms of the electrical property,
thermal stability and conductivity. The experimental results show that higher
GO concentrations, higher reduction temperatures and longer reduction times
increase the electrical conductivities of the GAs. The GAs synthesized with
more efficient reducing agents exhibit a higher electrical conductivity than
those with less efficient reducing agents. The addition of CNTs enhances the
electrical conductivities of pure GAs up to 2.8 times, while thermal annealing at
450 °C in Ar can lead to an increase of up to 5 times. The thermal stability of the
GAs is mainly affected by the morphological characteristics of the GAs. In
general, a more compacted structure leads to better thermal stability. The DTA
peak temperatures can possibly be optimized by adjusting the GO
concentrations, synthesis conditions, selection of reducing agents and CNT
inclusions.
97
A comparative infrared technique has been developed to measure the
thermal conductivities of the GAs; it reports the first benchmark results of the
thermal conductivities of GAs as a function of GO concentrations, and also
investigates the effects of thermal annealing on their thermal properties. For this
thesis, the thermal conductivity of the GAs is measured to be 0.12–0.36 W/m·K,
much lower than that of pristine graphene. These low values are likely due to
the high porosity of the GAs, the low quality and small size of the chemically
derived graphene, and the large thermal boundary resistance at the
graphene–graphene and graphene–air contacts.
In general, the thermal conductivity of the GAs can be controlled by
optimizing the GO concentration and other synthesis conditions (i.e. synthesis
temperature and time). Post-thermal annealing increases the thermal
conductivities of the GAs synthesized with a GO suspension of low graphene
concentration. In addition, by fitting the data with the rule of mixtures and
effective medium theory, the thermal conductivity of rGO is estimated to be
12.2–30.2 W/m·K. Due to their unique mesoporous network structure, GAs
have great potential to be developed into graphene-based composites by taking
advantage of capillary forces to infiltrate the porous network.
The multifunctional properties of the GAs will be summarized again in the
Conclusion chapter.
98
CHAPTER 6: Advanced Multifunctional Graphene
Aerogel–Poly (methyl methacrylate) Composites:
Experiments and Modelling
6.1 Introduction
Poly (methyl methacrylate) (PMMA) is a common type of thermoplastic.
Because of its unique advantages, such as optical clarity, low density, low cost,
and good physic-mechanical properties, PMMA has been widely used in the
medicine, optical, automobile, architectural, and construction industries.
Several articles have demonstrated the successful incorporation of graphene
nanosheets into the PMMA matrix. The results show that at only 1 wt. % of
graphene loading, the elastic modulus and ultimate tensile strength for PMMA
increase by 80% and 20% respectively [227]
, while at 2.7 vol. % loading, the
measured electrical conductivity of PMMA particles reaches 64 S/m [66]
. In
addition, at ~5 wt. % loading of graphene nanoplatelets, the glass transition
temperature Tg for PMMA is also reported to increase by 30% [278]
.
In this chapter, we develop a new approach to prepare highly conductive
and ultra-strong GA–PMMA nanocomposites by backfilling PMMA into GAs.
As presented in Section 3.2, the GAs were prepared through the modified
Hummers’ and chemical reduction method from graphite. PMMA was in situ
polymerized from the MMA monomer and infiltrated into the pores of the GAs
by capillary forces. The microhardness, electrical, and thermal properties of
these GA–PMMA nanocomposites were investigated. EMT theory was applied
to predict the thermal boundary resistance (TBR) of graphene–PMMA contact.
99
6.2 Structure of GA–PMMA nanocomposites
Upon chemical reduction, robust porous GAs as shown in Figure 6.1(a)
below are self-assembled by partially overlapped graphene nanosheets due to
π–π stacking. The GA–PMMA nanocomposites are obtained after the liquid
polymer precursor (pre-polymerized PMMA) is completely infused and cured
within the framework of the GAs. Figure 6.1(b) below shows an as-prepared
GA–PMMA nanocomposite sample polished to a square shape with
dimensions of 10 mm (L) × 10 mm (W) × 1.2 mm (H). Table 6-1 below
presents a summary of the synthesis conditions, graphene volume fraction and
multiple advanced properties of the GA–PMMA nanocomposites. The GAs are
synthesized from the GO concentrations of 1, 2, 6, and 12 mg/ml respectively,
and thus their corresponding graphene volume fractions are calculated based on
the density ratio of these as-prepared GAs to pristine graphene.
To further investigate its morphological features and evaluate the infusion
efficiency of PMMA into GAs, the cross-section of the GA–PMMA
nanocomposite is subjected to FE-SEM. As the image shown in Figure 6.1(d),
the GA network (see Figure 6.1(c) below) is completely infiltrated with PMMA
and the graphene nanosheets also adhere well to the PMMA. As with the
reported interaction between graphene and other polymers [45, 192, 227, 279]
, a good
wetting behavior is also considered to exist here between the reduced graphene
oxide (rGO) and PMMA. This can be potentially attributed to the formation of
hydrogen bonds between the remaining oxygen-containing groups of rGO and
the carbonyl of PMMA, which is analogous to the crystalline interfacial region
100
reported for poly (vinyl alcohol)-based composites.
Figure 6.1 Digital images of (a) GA and (b) GA–PMMA nanocomposite and
FE-SEM images of (c) the porous GA without PMMA infiltrated and (d) the
cross section of the GA–PMMA nanocomposite.
Table 6-1 Synthesis conditions, graphene volume fraction, and multifunctional
properties of GA–PMMA nanocomposites.
No.
GO
conc.
(mg/ml)
Volume
fraction
(%)*
Electrical
conductivity
(S/m)
Vickers
microhardness
(MPa)
Thermal
conductivity
(W/ (K m))
1 0 0 10-14
205.0 ± 7.8 0.20 ± 0.01
2 1 0.67 0.160 ± 0.003 303.6 ± 29.1 0.35 ± 0.01
3 2 1.34 0.215 ± 0.006 342.5 ± 28.4 0.46 ± 0.03
4 6 2.04 0.458 ± 0.009 386.8 ± 36.1 0.58 ± 0.04
5 12 2.50 0.859 ± 0.033 462.5 ± 42.5 0.70 ± 0.03
*The volume fraction is converted from the density ratio of the GAs to graphene,
as described in Table 5-3.
In addition, as there is no interfacial cracking observed in Figure 6.1(d),
the skeleton of GAs in the nanocomposite may still remain a good intact after
being backfilled with PMMA. Hence, as compared with graphene nanoplatelets,
the GAs act not only as an ideal uniformly dispersed reinforcement material, but
also as a porous molding material for the graphene-based nanocomposites.
101
6.3 Electrical property of GA–PMMA nanocomposites
In order to obtain equipotential surfaces during the electrical property test,
a thin layer of silver paste (<100 μm) is applied to both the top and bottom sides
of the GA–PMMA nanocomposites (10 mm (L) × 10 mm (W) × 1mm (H)).
Figure 6.2(a) below shows the electrical conductivity of the GA–PMMA
nanocomposites with different graphene loadings. When the graphene content
increases from 0.67 to 2.50 vol. %, the electrical conductivity of the
nanocomposites increases dramatically from 0.160 to 0.859 S/m, which
represents an improvement of more than 13 orders of magnitude compared with
that of pure PMMA (10-14
S/m [224, 233, 280]
). This significant increase in electrical
conductivity is likely a result of the increasing number of current pathways
which are generated as the graphene volume fraction increases. In addition,
even at only 0.67 vol. % graphene, the electrical conductivity of 0.160 S/m can
be a sufficient level for electromagnetic interference materials [236, 281-283]
. These
results may imply that the conductive network of the GAs is well preserved
within the GA–PMMA nanocomposites even at very low graphene loadings,
which also agrees with the observations in Figure 6.1(d).
The electrical conductivity of GA–PMMA nanocomposites could be
theoretically described further by the percolation theory [41, 43, 65, 236, 240, 278, 284]
using a power law relationship:
𝜎 = 𝜎𝑓[(𝜑 − 𝜑𝑐) (1 − 𝜑𝑐)⁄ ]𝑡 (6.1)
where σf and φ are the electrical conductivity and volume fraction of graphene
respectively, φc is the percolation threshold, and t is the ‘universal critical
102
exponent’.
Figure 6.2 (a) The electrical conductivity (black dots) of GA–PMMA
nanocomposites as a function of graphene volume fraction at room temperature.
The dashed blue line is the best fit from the power law relation 𝜎𝑐 =𝜎𝑓[(𝜑 − 𝜑𝑐) (1 − 𝜑𝑐)⁄ ]𝑡 with a nonlinear least-squares algorithm, where the
fitted values are σf=2543.8, φc=0 and t=2.18, respectively. (b) Comparative
study of electrical conductivity between this work and previously reported
graphene–PMMA nanocomposites synthesized via in situ polymerization. For
comparison purposes, the values of graphene loading reported in wt. % have
been converted to vol. %.
When the electrical conductivities of the GA–PMMA nanocomposites are
fitted with a nonlinear least-squares algorithm, the best agreement occurs at
103
φc=0 with the relation: 𝜎𝑐 = 2543.8𝜑2.18. It should be noted that the ‘zero
percolation threshold’ for the GA–PMMA nanocomposites is different from the
electrical behavior of the nanocomposites prepared from powdery graphene
fillers, but is consistent with that of other foam structures [285]
. In addition, the
electrical conductivity of rGO can be extracted to be 2543.8 S/m accordingly,
which is very close to the measured value of 2800 S/m reported by Zeng et al.
[224]
All the electrical conductivities measured in this work are approximately
one order greater than those of previously reported graphene–PMMA
composites (as shown in Figure 6.2(b) previously), indicating that the use of
GAs to develop graphene–polymer nanocomposites may be a more effective
way to solve the agglomeration problem of graphene nanosheets in the matrix.
6.4 Thermal stability of GA–PMMA nanocomposites
To investigate the thermal stability of the GA–PMMA nanocomposites, as
well as the extent of interaction between graphene nanosheets and PMMA, the
thermal degradations of pure PMMA and the GA–PMMA nanocomposite with
2.50 vol. % graphene loading are evaluated for comparison, as shown in Figure
6.3. Compare with the weight loss profile of pure PMMA, the TGA curve of the
GA–PMMA nanocomposite remarkably shifts towards a higher temperature.
The onset degradation temperature (the temperature at 5% weight loss) [286]
of
the GA–PMMA nanocomposite increases by 16 °C while the maximum
degradation temperature (the temperature at 50% weight loss) [231]
increases by
37 °C compared with pure PMMA. This significant enhancement in thermal
104
stability could also be attributed to the incorporation of homogeneously
dispersed graphene nanosheets, which act as barriers to inhibit the mobility of
PMMA chains during pyrolysis [63, 231-233]
.
Figure 6.3 Thermogravimetric curves of PMMA and GA–PMMA
nanocomposites.
Additionally, while pure PMMA completely degrades above ~380 °C, ~4
wt. % residual weight could still be observed in the GA–PMMA nanocomposite
until ~550 °C, due to the presence of thermally stable graphene nanosheets [44,
286], which is consistent with the converted weight fraction of graphene in the
GA–PMMA nanocomposite.
6.5 Thermal property of GA–PMMA nanocomposites
The results, plotted in Figure 6.4 below, are similar to those of electrical
conductivity. The thermal conductivity of the GA–PMMA nanocomposites also
increases with the increase in graphene concentration. The GA with 2.50 vol. %
graphene loading yields a thermal conductivity of as high as 0.70 W/m·K,
which is 3.5 times that of the pure PMMA (0.20 W/m·K) used here [287-289]
, and
105
also nearly 2 times greater than that of GO–filled PMMA nanocomposites [290]
.
Compared with the nanocomposites containing the same concentration but
non-uniformly dispersed graphene, the high thermal conductivity of the GA
network may provide much more effective channels for phonons to transfer.
Additionally, a strong interface and large contact area between the graphene and
the polymer matrix are also key factors that help increase the thermal
conductivity of the nanocomposites [248]
.
Figure 6.4 The thermal conductivity of GA–PMMA nanocomposites as a
function of graphene volume fraction. Inset is the set-up scheme for the
comparative infrared microscopy technique, where the amorphous quartz
having thermal conductivity of 1.3 W/m·K is used as the reference material. To
extract the thermal conductivity of rGO and TBR between rGO and PMMA, the
experimental data is fitted with a modified effective medium theory
simultaneously by a nonlinear least-squares algorithm. The red dashed line
represents the best fitting with the thermal conductivity of rGO (30.0 W/ m·K)
and the rGO-PMMA TBR (3.78×10-8
Km2/W).
However, considering the outstanding thermal conductivity of graphene
(5300 W/m·K [55]
), this improvement in the GA–PMMA nanocomposites seems
rather modest. This great mismatch is not only caused by the low intrinsic
thermal conductivity of rGO [143, 144]
, but also the large thermal resistance (also
106
known as Kapitza [249, 250, 291]
resistance) at the graphene–graphene [291]
and
graphene–PMMA interfaces [249]
.
A modified effective medium theory (EMT) is applied to model the
thermal conductivity of the GA–PMMA nanocomposites through the
relationship:
kGA−PMMA = krGO [3kPMMA+2f(krGO−kPMMA)
(3−f)krGO+kPMMAf+RBkPMMAkrGOf
H
] (6.2)
where RB is the TBR between rGO and PMMA, and H is the total thickness of
the rGO, which is taken to be 3 nm on average in this calculation [172, 276, 292]
.
When rGO nanosheets are assumed to be uniformly dispersed in PMMA and
not interconnected, an rGO thermal conductivity of ~30.0 W/m·K and a TBR
value of 3.78×10-8
km2/W are extracted from the best fit. The results are
consistent with our previous estimation for the thermal conductivity of rGO [292]
and on the same order of other reported TBR values for graphene–polymer
systems (~10-8
km2/W)
[247], respectively. Although the direct rGO–rGO
conduction is neglected in this model, the graphene-graphene boundary
resistance [(6.5–7 .7)×10-8
km2/W]
[291] is likely to be close to the TBR between
rGO and PMMA, and thus the disparity of our estimation may not be
significant.
6.6 Microhardness of GA–PMMA nanocomposites
In order to investigate the improvement in hardness, microindentation tests
are performed on the GA–PMMA nanocomposites and compared with pure
PMMA [282, 283]
. An increasing trend of microhardness with the increase of
graphene nanosheets can be clearly observed in Figure 6.5 below. With 0.67 to
107
2.50 vol. % graphene incorporated, the hardness values for the GA–PMMA
nanocomposites exhibit an increase from 48% to more than twice compared
with pure PMMA (215 MPa). In this work, the GA with 2.50 vol. % graphene
yields a maximum hardness of 462.5 MPa, which is significantly higher than
the first benchmark result (153.1 MPa at 0.6 wt. % graphene) reported by Das et
al. [293]
, and is even comparable to that of PMMA nanocomposites reinforced by
high-quality exfoliated graphite [281]
(as shown in Figure 6.5).
Besides the impact of the intrinsic high hardness and strength of graphene
[294], a strong interface between rGO and PMMA is considered an essential
reason for this great enhancement [293, 294]
. The presence of rGO nanosheets is
likely to generate numerous obstacles and interlock the motion of shear bands in
the polymer matrix, which might offer a better resistance to plastic deformation
and achieve more efficient load [227]
. Furthermore, the impact of the nucleation
of an rGO–PMMA interfacial region might also be significant, which is likely
to be much stiffer and may probably result in better stress transfer than the rest
of its amorphous bulk [295]
.
108
Figure 6.5 The microhardness of GA–PMMA nanocomposites as a function of
graphene volume fractions. Inset is a typical indentation for the Vickers
microhardness test. For all GA–PMMA nanocomposites, the diagonal
dimensions range from 25 to 60 μm. A prediction of the microhardness of
GA–PMMA nanocomposites, using the modified Halpin–Tsai equation, and a
comparative study of microhardness between this work and previously reported
graphene–PMMA nanocomposites are also included. For comparison purposes,
the values of graphene loading reported in mass fraction (wt. %) are converted
to volume fraction (vol. %).
Current results in terms of the mechanical property of graphene–polymer
nanocomposites appear to be scattered. Although different polymerization
methods are applied by different groups, our fabrication method still shows a
better capability in uniformly distributing graphene in composites, which could
109
potentially provide even higher microhardness levels for the GA–PMMA
nanocomposites at higher graphene loadings in the GAs.
The modified Halpin–Tsai equation [296-299]
, which has been widely used to
predict the mechanical properties of nanocomposites containing fiber or platelet
fillers, is applied to model the experimental results of the microhardness of
GA–PMMA nanocomposites in this work, as shown in Figure 6.5 above. For
randomly oriented nanoplatelets, the microhardness of GA–PMMA
nanocomposites containing φ (in vol. %) fillers can be estimated by Equation
(6.3) below:
𝐻c = 𝐻𝑚 (1+2𝛼𝜂𝜑
1−𝜂𝜑) (6.3)
where Hc and Hm are the hardness of the GA–PMMA nanocomposite and pure
PMMA respectively, α=l/d is the aspect ratio of graphene nanosheets, and η is
given by Equation (6.4) below:
𝜂 =𝐻𝑓−𝐻𝑚
𝐻𝑓+2𝛼𝐻𝑚 (6.4)
where Hf is the hardness for graphene nanosheets. When Hf and α are taken as
10.8 GPa and 400 respectively, the prediction in Figure 6.5 above shows an
excellent agreement with the experimental data. This theoretical analysis also
suggests that an effective load transfer occurs between the fillers and the matrix,
which can further confirm the good dispersion of graphene in PMMA and a
strong bond between the graphene and PMMA.
6.7 Conclusions
In summary, we have demonstrated a novel method of fabricating
GA–PMMA nanocomposites by infiltrating a PMMA precursor into the
110
network of the GAs, which is followed by curing. Because of the uniform
distribution of graphene nanosheets in the PMMA matrix, when the graphene
loading increases from 0.67 to 2.50 vol. %, the GA–PMMA nanocomposites
exhibit significant enhancements in their electrical conductivities (0.160–0.859
S/m), microhardness (303.6–462.5 MPa), and thermal conductivities
(0.35–0.70 W/m·K), compared with pure PMMA and graphene–PMMA
nanocomposites prepared by traditional powdery dispersion methods. The
percolation theory, Halpin–Tsai equation, and effective medium theory are
applied to model the electrical, microhardness, and thermal properties,
respectively, of the GA–PMMA nanocomposites, and all yield good
agreements with the experimentally measured data. The thermal boundary
resistance of the graphene–PMMA is estimated to be 3.78×10-8
Km2/W in this
work. This novel fabrication method for the GA–PMMA nanocomposites may
effectively solve the dispersion issue of graphene (or CNTs) in the polymer
matrix, and also offer a facile approach for the preparation of various
filler-matrix nanocomposites with multifunctional capabilities. These
GA–PMMA are expected to find diverse applications in the fields of aerospace,
tissue scaffolds [43]
, composite electrodes [43]
, biosensors [300]
, energy generation,
and storage [301]
.
The development and multifunctional properties of the GA–PMMA
nanocomposites will be summarized again in the Conclusion chapter.
111
CHAPTER 7: Conclusions and Recommendations
7.1 Conclusions
Graphene aerogels (GAs), having large surface area property of aerogels
and the excellent multifunctional properties of graphene nanosheets, can be
promising candidates for several applications, such as energy storage devices
and light-weight nanocomposites with advanced multifunctional properties.
From the findings of this thesis, the following conclusions can be drawn.
Firstly, this study has developed a cost-effective method to fabricate
self-assembled GAs from graphite at a low temperature range (45–95 oC). The
effects of chemical composition (initial GO concentration) and fabrication
conditions (such as reductions in temperature and time, reducing agents,
carbon nanotubes (CNTs), and thermal annealing processes) on the
morphology control, electrical property and thermal stability of the
as-prepared GAs have also been investigated comprehensively. It is found that
the GAs synthesized from a higher GO concentration have a higher surface area
and electrical conductivity, while higher reduction temperatures or longer
reduction times cause GAs to have a more packed structure with lower surface
area, higher electrical conductivity, and better thermal stability.
Similarly, the GAs synthesized with a more efficient reducing agent
exhibit a higher electrical conductivity and lower surface area, while those with
a low-efficiency reducing agent possess a lower electrical conductivity and
higher surface area. In addition, a thermal annealing treatment at 450 oC in
Argon is conducted on these GAs, further enhancing their electrical
112
conductivities by approximately 5 times. To explore the effects of CNT
inclusions on the morphology and multifunctional properties of GAs,
graphene–CNT hybrid aerogels are successfully developed by adding CNTs
during the GA synthesis process. CNTs can increase the surface area of GAs up
to 14% and enhance their electrical conductivity by up to 2.8 times. The thermal
stability of the GAs is also affected by the choice of reducing agents and CNT
inclusions. An increased thermal stability is found in the GAs synthesized with
a high-efficiency reducing agent, while a decrease is found in those with a
dispersed CNT concentration above 1.0 mg/ml.
The experimental results of this thesis indicate that the morphology of GAs
can be effectively controlled by changing the synthesis parameters
appropriately; this also significantly impacts the multifunctional properties of
GAs in turn. This systematical study may provide insight into optimizing the
nanostructures and multifunctional properties of the GAs for various potential
applications, such as the development of electrode materials, energy storage
devices, and advanced nanocomposites in general.
Secondly, a comparative infrared technique to measure the thermal
conductivity and investigate the thermal conduction mechanisms of GAs has
been developed for the first time. This thesis reports the first benchmark results
of the thermal conductivity of GAs as a function of GO concentrations, and also
studies the effects of the thermal annealing process. The thermal conductivity of
the GAs with 0.67–2.50 vol. % graphene is measured to be 0.12–0.36 W/m·K,
which is several orders lower than that of pristine graphene. These low values
can likely be attributed to the high porosity of the GAs, the low quality and
113
small size of the chemically derived graphene, and the large thermal boundary
resistance at the graphene–graphene and graphene–air contacts. The thermal
conductivity of the GAs can be controlled by optimizing the GO concentration
and other synthesis conditions (i.e. synthesis temperature and time), while a
post-thermal annealing of the as-prepared GAs is found to be able to increase
the thermal conductivity of the GAs synthesized from the GO suspension of low
graphene concentration. Based on the measured thermal conductivity of the
GAs, existing models are also applied to predict the thermal conductivity of the
reduced graphene oxide (rGO) in the GAs. By fitting the data with rule of
mixtures and effective medium theory, the thermal conductivity of rGO is
estimated to be 12.2–30.2 W/m·K. This appears to be the highest value reported
so far for the thermal conductivity of rGO. In addition, the developed
comparative infrared technique can be further applied to measure the thermal
conductivity of various other porous materials.
Owing to their unique mesoporous network structure, GAs have great
potential to be developed into graphene-based nanocomposites by taking
advantage of capillary forces to infiltrate their porous network. Finally, a novel
method to fabricate GA–PMMA nanocomposites has been demonstrated, in
which the network of the GAs is infiltrated by a PMMA precursor and cured in
situ. When graphene loading is varied from 0.67 to 2.50 vol. %, the GA–PMMA
nanocomposites exhibit significant enhancements in electrical conductivity
(0.16–0.86 S/m), microhardness (303.6–462.5 MPa) and thermal conductivity
(0.35–0.70 W/m·K), compared with pure PMMA and the graphene–PMMA
nanocomposites prepared by traditional dispersion methods. This could be
114
attributed to the uniform distribution of the graphene nanosheets in the PMMA
matrix. The percolation theory, Halpin–Tsai equation, and effective medium
theory are applied to model the electrical, microhardness, and thermal
properties, respectively, of the GA–PMMA nanocomposites, and all yield good
agreements with the experimental data. The thermal boundary resistance of
graphene–PMMA is estimated to be 3.78×10-8
km2/W in this thesis. This novel
method of fabricating the GA–PMMA nanocomposites may effectively
overcome the dispersion challenge of graphene (or CNTs) in the polymer
matrix, thus offering an advanced approach for the preparation of various
filler-matrix nanocomposites with multifunctional capabilities.
7.2 Recommendations
However, it should be noted that for this thesis, only the chemical
reduction method is applied to synthesize GAs. The development of more
versatile and efficient methods of fabricating GAs and other carbon-based
aerogels with superior multifunctional properties will also be interesting to
investigate.
Moreover, the GAs are only applied to synthesize GA–PMMA
nanocomposites, while their application for energy storage devices has not been
investigated yet. Due to their high surface area and good multifunctional
properties, the morphology–controlled GAs may find promising applications in
batteries, fuel cells, and high-performance supercapacitors.
In addition, thermal boundary resistances (TBRs) at graphene–graphene
(Chapter 4) and graphene–PMMA (Chapter 5) interfaces are only estimated
115
through the modified EMT theory, in which the size distribution, quantities, and
dispersions of graphene cannot be well considered, and only the linear relation
between the effective thermal conductivity and graphene volume fraction is
possibly involved. Therefore, more accurate models (or approaches) that can
take the complex geometry of graphene sheets into account to predict the
thermal conductivity of aerogels and composites are especially needed.
116
REFERENCES
[1] A. Du, B. Zhou, Z. Zhang, J. Shen, A special material or a new state of
matter: A review and reconsideration of the aerogel, Materials, 2013, 6,
941-968.
[2] J. Fricke, A. Emmerling, Aerogels–Recent progress in production
techniques and novel applications, J. Sol-Gel Sci. Technol., 1998, 13, 299-303.
[3] I. M. El-Nahhal, N. M. El-Ashgar, A review on polysiloxane–immobilized
ligand systems: Synthesis, characterization and applications, J. Organomet.
Chem., 2007, 692, 2861-2886.
[4] H. Sun, Z. Xu, C. Gao, Multifunctional, ultra-flyweight, synergistically
assembled carbon aerogels, Adv. Mater., 2013, 25, 2554-2560.
[5] A. J. Hunt, C. A. Jantzen, W. Q. Cao, Aerogel–A high performance
insulating material at 0.1 bar, Amer Soc Testing and Materials, 1991, pp.
455-463.
[6] A. E. Gash, T. M. Tillotson, J. H. Satcher Jr, L. W. Hrubesh, R. L. Simpson,
New sol–gel synthetic route to transition and main-group metal oxide aerogels
using inorganic salt precursors, J. Non-Cryst. Solids, 2001, 285, 22-28.
[7] D. J. Suh, T. J. Park, Sol-gel strategies for pore size control of
high-surface-area transition-metal oxide aerogels, Chem. Mater., 1996, 8,
509-513.
117
[8] R. Pekala, Organic aerogels from the polycondensation of resorcinol with
formaldehyde, J. Mater. Sci., 1989, 24, 3221-3227.
[9] Y. Guo, W. Meyer-Zaika, M. Muhler, S. Vukojević, M. Epple, Cu/Zn/Al
xerogels and aerogels prepared by a sol–gel reaction as catalysts for methanol
synthesis, Eur. J. Inorg. Chem., 2006, 2006, 4774-4781.
[10] R. Pekala, J. Farmer, C. Alviso, T. Tran, S. Mayer, J. Miller, B. Dunn,
Carbon aerogels for electrochemical applications, J. Non-Cryst. Solids, 1998,
225, 74-80.
[11] M. B. Bryning, D. E. Milkie, M. F. Islam, L. A. Hough, J. M. Kikkawa, A.
G. Yodh, Carbon nanotube aerogels, Adv. Mater., 2007, 19, 661-664.
[12] R. R. Kohlmeyer, M. Lor, J. Deng, H. Liu, J. Chen, Preparation of stable
carbon nanotube aerogels with high electrical conductivity and porosity,
Carbon, 2011, 49, 2352-2361.
[13] J. Li, J. Li, H. Meng, S. Xie, B. Zhang, L. Li, H. Ma, J. Zhang, M. Yu,
Ultra-light, compressible and fire-resistant graphene aerogel as a highly
efficient and recyclable absorbent for organic liquids, J. Mater. Chem. A, 2014,
2, 2934-2941.
[14] X. Zhang, Z. Sui, B. Xu, S. Yue, Y. Luo, W. Zhan, B. Liu, Mechanically
strong and highly conductive graphene aerogel and its use as electrodes for
electrochemical power sources, J. Mater. Chem., 2011, 21, 6494-6497.
118
[15] S. Nardecchia, D. Carriazo, M. L. Ferrer, M. C. Gutierrez, F. del Monte,
Three dimensional macroporous architectures and aerogels built of carbon
nanotubes and/or graphene: Synthesis and applications, Chem. Soc. Rev., 2013,
42, 794-830.
[16] L. Wang, M. Wang, Z. Huang, T. Cui, X. Gui, F. Kang, K. Wang, D. Wu,
Capacitive deionization of NaCl solutions using carbon nanotube sponge
electrodes, J. Mater. Chem., 2011, 21, 18295-18299.
[17] J. Biener, M. Stadermann, M. Suss, M. A. Worsley, M. M. Biener, K. A.
Rose, T. F. Baumann, Advanced carbon aerogels for energy applications,
Energy Environ. Sci., 2011, 4, 656-667.
[18] L. W. Hrubesh, Aerogel applications, J. Non-Cryst. Solids, 1998, 225,
335-342.
[19] F. Meng, X. Zhang, B. Xu, S. Yue, H. Guo, Y. Luo, Alkali-treated
graphene oxide as a solid base catalyst: Synthesis and electrochemical
capacitance of graphene/carbon composite aerogels, J. Mater. Chem., 2011, 21,
18537-18539.
[20] Z. Wu, S. Yang, Y. Sun, K. Parvez, X. Feng, K. Müllen, 3D nitrogen-doped
graphene aerogel-supported Fe3O4 nanoparticles as efficient electrocatalysts
for the oxygen reduction reaction, J. Am. Chem. Soc., 2012, 134, 9082-9085.
[21] S. Berber, Y. Kwon, D. Tománek, Unusually high thermal conductivity of
carbon nanotubes, Phys. Rev. Lett., 2000, 84, 4613-4616.
119
[22] A. M. Cassell, J. A. Raymakers, J. Kong, H. Dai, Large scale CVD
synthesis of single-walled carbon nanotubes, J. Phys. Chem. B, 1999, 103,
6484-6492.
[23] W. Lu, M. Zu, J. H. Byun, B. S. Kim, T. W. Chou, State of the art of carbon
nanotube fibers: Opportunities and challenges, Adv. Mater., 2012, 24,
1805-1833.
[24] A. K. Geim, K. S. Novoselov, The rise of graphene, Nat. Mater., 2007, 6,
183-191.
[25] H. Zhou, W. Yao, G. Li, J. Wang, Y. Lu,
Graphene/poly(3,4-ethylenedioxythiophene) hydrogel with excellent
mechanical performance and high conductivity, Carbon, 2013, 59, 495-502.
[26] S. Ansari, E. P. Giannelis, Functionalized graphene sheet–poly(vinylidene
fluoride) conductive nanocomposites, J. Polym. Sci., Part B: Polym. Phys.,
2009, 47, 888-897.
[27] C. A. García-González, M. C. Camino-Rey, M. Alnaief, C. Zetzl, I.
Smirnova, Supercritical drying of aerogels using CO2: Effect of extraction time
on the end material textural properties, J. Supercrit. Fluids., 2012, 66, 297-306.
[28]G. W. Scherer, D. M. Smith, Cavitation during drying of a gel, J. Non-Cryst.
Solids, 1995, 189, 197-211.
120
[29] H. Namatsu, K. Yamazaki, K. Kurihara, Supercritical drying for
nanostructure fabrication without pattern collapse, Microelectron. Eng., 1999,
46, 129-132.
[30] F. Di Renzo, R. Valentin, M. Boissière, A. Tourrette, G. Sparapano, K.
Molvinger, J. M. Devoisselle, C. Gérardin, F. Quignard, Hierarchical
macroporosity induced by constrained syneresis in core−shell polysaccharide
composites, Chem. Mater., 2005, 17, 4693-4699.
[31] D. T. Carty, T. G. Dewees, R. J. Iliff, F. M. Knafelc, J. R. Latham, T. M.
Lipton, J. D. Mitchell, R. G. Taylor, Liquid/supercritical carbon dioxide dry
cleaning system, Google Patents, 1993.
[32] B. M. Novak, D. Auerbach, C. Verrier, Low-density, mutually
interpenetrating organic-inorganic composite materials via supercritical drying
techniques, Chem. Mater., 1994, 6, 282-286.
[33] M. B. Bryning, Carbon nanotube networks in epoxy composites and
aerogels, Physics and Astronomy, University of Pennsylvania, 2007.
[34] F. Franks, Freeze-drying of bioproducts: Putting principles into practice,
Eur. J. Pharm. Biopharm., 1998, 45, 221-229.
[35] H. Kabbour, T. F. Baumann, J. H. Satcher, A. Saulnier, C. C. Ahn, Toward
new candidates for hydrogen storage: High-surface-area carbon aerogels,
Chem. Mater., 2006, 18, 6085-6087.
121
[36] A. C. Pierre, G. M. Pajonk, Chemistry of aerogels and their applications,
Chem. Rev., 2002, 102, 4243-4266.
[37] L. Forest, V. Gibiat, T. Woignier, Evolution of the acoustical properties of
silica alcogels during their formation, Ultrasonics, 1998, 36, 477-481.
[38] M. Schmidt, F. Schwertfeger, Applications for silica aerogel products, J.
Non-Cryst. Solids, 1998, 225, 364-368.
[39] J. P. Randall, M. A. B. Meador, S. C. Jana, Tailoring mechanical properties
of aerogels for aerospace applications, ACS Appl. Mater. Interfaces, 2011, 3,
613-626.
[40] P. H. C. Camargo, K. G. Satyanarayana, F. Wypych, Nanocomposites:
Synthesis, structure, properties and new application opportunities, Mater. Res.,
2009, 12, 1-39.
[41] G. Gonçalves, P. A. A. P. Marques, A. Barros-Timmons, I. Bdkin, M. K.
Singh, N. Emami, J. Grácio, Graphene oxide modified with PMMA via ATRP
as a reinforcement filler, J. Mater. Chem., 2010, 20, 9927-9934.
[42] J. Loomis, X. Fan, F. Khosravi, P. Xu, M. Fletcher, R. W. Cohn, B.
Panchapakesan, Graphene/elastomer composite-based photo-thermal
nanopositioners, Sci. Rep., 2013, 3, 1900.
[43] J. R. Potts, D. R. Dreyer, C. W. Bielawski, R. S. Ruoff, Graphene-based
polymer nanocomposites, Polymer, 2011, 52, 5-25.
122
[44] K. Zhang, W. L. Zhang, H. J. Choi, Facile fabrication of self-assembled
PMMA/graphene oxide composite particles and their electroresponsive
properties, Colloid. Polym. Sci., 2012, 291, 955-962.
[45] R. Verdejo, M. M. Bernal, L. J. Romasanta, M. A. Lopez-Manchado,
Graphene filled polymer nanocomposites, J. Mater. Chem., 2011, 21,
3301-3310.
[46] V. Singh, D. Joung, L. Zhai, S. Das, S. I. Khondaker, S. Seal, Graphene
based materials: Past, present and future, Prog. Mater. Sci., 2011, 56,
1178-1271.
[47] J. Du, H. Cheng, The fabrication, properties, and uses of graphene/polymer
composites, Macromol. Chem. Phys., 2012, 213, 1060-1077.
[48] F. Irin, S. Das, F. O. Atore, M. J. Green, Ultralow percolation threshold in
aerogel and cryogel templated composites, Langmuir, 2013, 29, 11449-11456.
[49] K. S. Novoselov, A. K. Geim, S. V. Morozov, D. Jiang, Y. Zhang, S. V.
Dubonos, I. V. Grigorieva, A. A. Firsov, Electric field effect in atomically thin
carbon films, Science, 2004, 306, 666-669.
[50]P. Blake, P. D. Brimicombe, R. R. Nair, T. J. Booth, D. Jiang, F. Schedin, L.
A. Ponomarenko, S. V. Morozov, H. F. Gleeson, E. W. Hill, A. K. Geim, K. S.
Novoselov, Graphene-based liquid crystal device, Nano Lett., 2008, 8,
1704-1708.
123
[51] H. K. Chae, D. Y. Siberio-Perez, J. Kim, Y. Go, M. Eddaoudi, A. J.
Matzger, M. O. Keeffe, O. M. Yaghi, A route to high surface area, porosity and
inclusion of large molecules in crystals, Nature, 2004, 427, 523-527.
[52] Y. Zhu, S. Murali, W. Cai, X. Li, J. W. Suk, J. R. Potts, R. S. Ruoff,
Graphene and graphene oxide: Synthesis, properties, and applications, Adv.
Mater., 2010, 22, 3906-3924.
[53] K. S. Novoselov, A. K. Geim, S. V. Morozov, D. Jiang, M. I. Katsnelson, I.
V. Grigorieva, S. V. Dubonos, A. A. Firsov, Two-dimensional gas of massless
Dirac fermions in graphene, Nature, 2005, 438, 197-200.
[54] C. Lee, X. Wei, J. W. Kysar, J. Hone, Measurement of the elastic properties
and intrinsic strength of monolayer graphene, Science, 2008, 321, 385-388.
[55] A. A. Balandin, S. Ghosh, W. Bao, I. Calizo, D. Teweldebrhan, F. Miao, C.
N. Lau, Superior thermal conductivity of single-layer graphene, Nano Lett.,
2008, 8, 902-907.
[56] H. Chang, Z. Sun, Q. Yuan, F. Ding, X. Tao, F. Yan, Z. Zheng, Thin film
field-effect phototransistors from bandgap-tunable, solution-processed,
few-layer reduced graphene oxide films, Adv. Mater., 2010, 22, 4872-4876.
[57]Q. Yan, B. Huang, J. Yu, F. Zheng, J. Zhang, J. Wu, B. Gu, F. Liu, W. Duan,
Intrinsic current–voltage characteristics of graphene nanoribbon transistors and
effect of edge doping, Nano Lett., 2007, 7, 1469-1473.
124
[58] F. Xia, D. B. Farmer, Y. M. Lin, P. Avouris, Graphene field-effect
transistors with high on/off current ratio and large transport band gap at room
temperature, Nano Lett., 2010, 10, 715-718.
[59] S. De, P. J. King, M. Lotya, A. O'Neill, E. M. Doherty, Y. Hernandez, G. S.
Duesberg, J. N. Coleman, Flexible, transparent, conducting films of randomly
stacked graphene from surfactant-stabilized, oxide-free graphene dispersions,
Small, 2010, 6, 458-464.
[60] M. D. Stoller, S. Park, Y. Zhu, J. An, R. S. Ruoff, Graphene-based
ultracapacitors, Nano Lett., 2008, 8, 3498-3502.
[61] Q. Wu, Y. Xu, Z. Yao, A. Liu, G. Shi, Supercapacitors based on flexible
graphene/polyaniline nanofiber composite films, ACS Nano, 2010, 4,
1963-1970.
[62] A. M. Marconnet, N. Yamamoto, M. A. Panzer, B. L. Wardle, K. E.
Goodson, Thermal conductivity in aligned carbon nanotube-polymer
nanocomposites with high packing density, ACS Nano, 2011, 5, 4818-4825.
[63] J. Wang, H. Hu, X. Wang, C. Xu, M. Zhang, X. Shang, Preparation and
mechanical and electrical properties of graphene nanosheets–poly(methyl
methacrylate) nanocomposites via in situ suspension polymerization, J. Appl.
Polym. Sci., 2011, 122, 1866-1871.
[64] L. Dai, D. W. Chang, J. B. Baek, W. Lu, Carbon nanomaterials for
advanced energy conversion and storage, Small, 2012, 8, 1130-1166.
125
[65] X. Huang, X. Qi, F. Boey, H. Zhang, Graphene-based composites, Chem.
Soc. Rev., 2012, 41, 666-686.
[66] V. H. Pham, T. T. Dang, S. H. Hur, E. J. Kim, J. S. Chung, Highly
conductive poly(methyl methacrylate) (PMMA)–reduced graphene oxide
composite prepared by self-assembly of PMMA latex and graphene oxide
through electrostatic interaction, ACS Appl. Mater. Interfaces, 2012, 4,
2630-2636.
[67] J. Slonczewski, P. Weiss, Band structure of graphite, Phy. Rev., 1958, 109,
272-279.
[68] C. Soldano, A. Mahmood, E. Dujardin, Production, properties and potential
of graphene, Carbon, 2010, 48, 2127-2150.
[69] K. T. He, Characterization of graphene-substrate interactions using
scanning tunneling microscopy, Electrical and Computer Engineering,
University of Illinois at Urbana-Champaign, 2013.
[70] P. Avouris, Graphene: Electronic and photonic properties and devices,
Nano Lett., 2010, 10, 4285-4294.
[71] A. H. C. Neto, F. Guinea, N. M. R. Peres, K. S. Novoselov, A. K. Geim,
The electronic properties of graphene, Rev. Mod. Phys., 2009, 81, 109-162.
[72] T. A. Nguyen, J. U. Lee, D. Yoon, H. Cheong, Excitation energy dependent
Raman signatures of ABA- and ABC-stacked few-layer graphene, Sci. Rep.,
2014, 4, 4630.
126
[73] Y. Shibuta, J. A. Elliott, Interaction between two graphene sheets with a
turbostratic orientational relationship, Chem. Phys. Lett., 2011, 512, 146-150.
[74] Y. Zhang, J. P. Small, W. V. Pontius, P. Kim, Fabrication and
electric-field-dependent transport measurements of mesoscopic graphite
devices, Appl. Phys. Lett., 2005, 86, 073104.
[75] X. Lu, M. Yu, H. Huang, R. R. S, Tailoring graphite with the goal of
achieving single sheets, Nanotechnology, 1999, 10, 269-272.
[76] X. Lu, H. Huang, N. Nemchuk, R. S. Ruoff, Patterning of highly oriented
pyrolytic graphite by oxygen plasma etching, Appl. Phys. Lett., 1999, 75,
193-195.
[77] K. S. Novoselov, D. Jiang, F. Schedin, T. J. Booth, V. V. Khotkevich, S. V.
Morozov, A. K. Geim, T. M. Rice, Two-dimensional atomic crystals, Proc.
Natl. Acad. Sci. U. S. A., 2005, 102, 10451-10453.
[78] V. E. Aleksandr, M. I. Inna, A. K. Andrei, N. K. Dmitrii, Graphene:
Fabrication methods and thermophysical properties, Phys. Usp., 2011, 54,
227-258.
[79]P. Blake, E. W. Hill, A. H. Castro Neto, K. S. Novoselov, D. Jiang, R. Yang,
T. J. Booth, A. K. Geim, Making graphene visible, Appl. Phys. Lett., 2007, 91,
063124.
[80] L. Gao, W. Ren, F. Li, H. Cheng, Total color difference for rapid and
accurate identification of graphene, ACS Nano, 2008, 2, 1625-1633.
127
[81] I. Jung, M. Pelton, R. Piner, D. A. Dikin, S. Stankovich, S. Watcharotone,
M. Hausner, R. S. Ruoff, Simple approach for high-contrast optical imaging
and characterization of graphene-based sheets, Nano Lett., 2007, 7, 3569-3575.
[82] L. M. Malard, M. A. Pimenta, G. Dresselhaus, M. S. Dresselhaus, Raman
spectroscopy in graphene, Phys. Rep., 2009, 473, 51-87.
[83] J. H. Chen, C. Jang, S. Xiao, M. Ishigami, M. S. Fuhrer, Intrinsic and
extrinsic performance limits of graphene devices on SiO2, Nat. Nanotechnol.,
2008, 3, 206-209.
[84] K. I. Bolotin, K. J. Sikes, Z. Jiang, M. Klima, G. Fudenberg, J. Hone, P.
Kim, H. L. Stormer, Ultrahigh electron mobility in suspended graphene, Solid
State Commun., 2008, 146, 351-355.
[85] M. S. Dresselhaus, G. Dresselhaus, Intercalation compounds of graphite,
Adv. Phys., 2002, 51, 1-186.
[86] Y. Hernandez, V. Nicolosi, M. Lotya, F. M. Blighe, Z. Sun, S. De, I. T.
McGovern, B. Holland, M. Byrne, Y. K. Gun'Ko, J. J. Boland, P. Niraj, G.
Duesberg, S. Krishnamurthy, R. Goodhue, J. Hutchison, V. Scardaci, A. C.
Ferrari, J. N. Coleman, High-yield production of graphene by liquid-phase
exfoliation of graphite, Nat. Nanotechnol., 2008, 3, 563-568.
[87] X. Li, G. Zhang, X. Bai, X. Sun, X. Wang, E. Wang, H. Dai, Highly
conducting graphene sheets and Langmuir-Blodgett films, Nat. Nanotechnol.,
2008, 3, 538-542.
128
[88] J. W. May, Platinum surface LEED rings, Surf. Sci., 1969, 17, 267-270.
[89] J. C. Shelton, H. R. Patil, J. M. Blakely, Equilibrium segregation of carbon
to a nickel(111) surface: A surface phase transition, Surf. Sci., 1974, 43,
493-520.
[90] M. Eizenberg, J. M. Blakely, Carbon interaction with nickel surfaces:
Monolayer formation and structural stability, J. Chem. Phys., 1979, 71, 3467.
[91] M. Eizenberg, J. M. Blakely, Carbon monolayer phase condensation on
Ni(111), Surf. Sci., 1979, 82, 228-236.
[92] J. C. Hamilton, J. M. Blakely, Carbon segregation to single crystal surfaces
of Pt, Pd and Co, Surf. Sci., 1980, 91, 199-217.
[93] T. Fujita, W. Kobayashi, C. Oshima, Novel structures of carbon layers on a
Pt(111) surface, Surf. Interface Anal., 2005, 37, 120-123.
[94]A. Charrier, A. Coati, T. Argunova, F. Thibaudau, Y. Garreau, R. Pinchaux,
I. Forbeaux, J. M. Debever, M. Sauvage-Simkin, J. M. Themlin, Solid-state
decomposition of silicon carbide for growing ultra-thin heteroepitaxial graphite
films, J. Appl. Phys., 2002, 92, 2479.
[95] L. Papagno, L. Caputi, Determination of graphitic carbon structure
adsorbed on Ni(110) by surface extended energy-loss fine-structure analysis,
Phys. Rev. B, 1984, 29, 1483-1486.
129
[96] T. Tanaka, A. Itoh, K. Yamashita, E. Rokuta, C. Oshima, Heteroepitaxial
system of h–BN/monolayer graphene on Ni(111), Surf. Rev. Lett., 2003, 10,
697-703.
[97] C. Berger, Z. Song, T. Li, X. Li, A. Y. Ogbazghi, R. Feng, Z. Dai, A. N.
Marchenkov, E. H. Conrad, P. N. First, W. A. d. Heer, Ultrathin epitaxial
graphite: 2D electron gas properties and a route toward graphene-based
nanoelectronics, J. Phys. Chem. B, 2004, 108, 19912-19916.
[98] C. Berger, Z. Song, X. Li, X. Wu, N. Brown, C. Naud, D. Mayou, T. Li, J.
Hass, A. N. Marchenkov, E. H. Conrad, P. N. First, W. A. d. Heer, Electronic
confinement and coherence in patterned epitaxial graphene, Science, 2006, 312,
1191-1196.
[99] T. Ohta, B. Aaron, T. Seyller, K. Horn, E. Rotenberg, Controlling the
electronic structure of bilayer graphene, Science, 2006, 313, 951-954.
[100] A. J. Van Bommel, J. E. Crombeen, A. Van Tooren, LEED and Auger
electron observations of the SiC(0001) surface, Surf. Sci., 1975, 48, 463-472.
[101] W. A. de Heer, C. Berger, X. Wu, P. N. First, E. H. Conrad, X. Li, T. Li,
M. Sprinkle, J. Hass, M. L. Sadowski, M. Potemski, G. Martinez, Epitaxial
graphene, Solid State Commun., 2007, 143, 92-100.
[102] C. Faugeras, A. Nerriere, M. Potemski, A. Mahmood, E. Dujardin, C.
Berger, W. A. de Heer, Few-layer graphene on SiC, pyrolitic graphite, and
graphene: A Raman scattering study, Appl. Phys. Lett., 2008, 92, 011914.
130
[103] J. Hass, F. Varchon, J. Millán-Otoya, M. Sprinkle, N. Sharma, W. de
Heer, C. Berger, P. First, L. Magaud, E. Conrad, Why multilayer graphene on
4H-SiC(0001) behaves like a single sheet of graphene, Phys. Rev. Lett., 2008,
100.
[104] K. V. Emtsev, A. Bostwick, K. Horn, J. Jobst, G. L. Kellogg, L. Ley, J.
L. McChesney, T. Ohta, S. A. Reshanov, J. Rohrl, E. Rotenberg, A. K. Schmid,
D. Waldmann, H. B. Weber, T. Seyller, Towards wafer-size graphene layers by
atmospheric pressure graphitization of silicon carbide, Nat. Mater., 2009, 8,
203-207.
[105] M. C. Lemme, T. J. Echtermeyer, M. Baus, H. Kurz, A graphene
field-effect device, Electron Device Letters, IEEE, 2007, 28, 282-284.
[106] Y. Q. Wu, P. D. Ye, M. A. Capano, Y. Xuan, Y. Sui, M. Qi, J. A.
Cooper, T. Shen, D. Pandey, G. Prakash, R. Reifenberger, Top-gated graphene
field-effect-transistors formed by decomposition of SiC, Appl. Phys. Lett., 2008,
92, 092102.
[107] T. A. Land, T. Michely, R. J. Behm, J. C. Hemminger, G. Comsa, STM
investigation of single layer graphite structures produced on Pt(111) by
hydrocarbon decomposition, Surf. Sci., 1992, 264, 261-270.
[108] P. W. Sutter, J. I. Flege, E. A. Sutter, Epitaxial graphene on ruthenium,
Nat. Mater., 2008, 7, 406-411.
131
[109] W. J. Arnoult, R. B. McLellan, The solubility of carbon in rhodium
ruthenium, iridium and rhenium, Scr. Metall., 1972, 6, 1013-1018.
[110] F. Varchon, R. Feng, J. Hass, X. Li, B. Nguyen, C. Naud, P. Mallet, J.
Y. Veuillen, C. Berger, E. Conrad, L. Magaud, Electronic structure of epitaxial
graphene layers on SiC: Effect of the substrate, Phys. Rev. Lett., 2007, 99,
126805.
[111] K. S. Kim, Y. Zhao, H. Jang, S. Y. Lee, J. M. Kim, J. H. Ahn, P. Kim,
J. Y. Choi, B. H. Hong, Large-scale pattern growth of graphene films for
stretchable transparent electrodes, Nature, 2009, 457, 706-710.
[112] A. Reina, X. Jia, J. Ho, D. Nezich, H. Son, V. Bulovic, M. S.
Dresselhaus, J. Kong, Large area, few-layer graphene films on arbitrary
substrates by chemical vapor deposition, Nano Lett., 2008, 9, 30-35.
[113] X. Li, W. Cai, J. An, S. Kim, J. Nah, Large-area synthesis of
high-quality and uniform graphene films on copper foils, Science (New York,
N.Y.), 2009, 324, 1312-1314.
[114] Y. Si, E. T. Samulski, Synthesis of water soluble graphene, Nano Lett.,
2008, 8, 1679-1682.
[115] S. Park, R. S. Ruoff, Chemical methods for the production of
graphenes, Nat. Nanotechnol., 2009, 4, 217-224.
[116] B. C. Brodie, On the atomic weight of graphite, Phil. Trans. R. Soc.
Lond., 1859, 149, 249-259.
132
[117] W. S. Hummers, R. E. Offeman, Preparation of graphitic oxide, J. Am.
Chem. Soc., 1958, 80, 1339.
[118] S. Duquesne, M. L. Bras, S. Bourbigot, R. Delobel, H. Vezin, G.
Camino, B. Eling, C. Lindsay, T. Roels, Expandable graphite: A fire retardant
additive for polyurethane coatings, Fire Mater., 2003, 27, 103-117.
[119] Y. Zhu, S. Murali, M. D. Stoller, A. Velamakanni, R. D. Piner, R. S.
Ruoff, Microwave assisted exfoliation and reduction of graphite oxide for
ultracapacitors, Carbon, 2010, 48, 2118-2122.
[120] H. P. Boehm, A. Clauss, U. Hofmann, G. O. Fischer, Dunnste
kohlenstoff–Folien, Z. Naturforsch., B: Chem. Sci., 1962, B 17, 150-153.
[121] N. I. Kovtyukhova, P. J. Ollivier, B. R. Martin, T. E. Mallouk, S. A.
Chizhik, E. V. Buzaneva, A. D. Gorchinskiy, Layer-by-layer assembly of
ultrathin composite films from micron-sized graphite oxide sheets and
polycations, Chem. Mater., 1999, 11, 771-778.
[122] J. L. Vickery, A. J. Patil, S. Mann, Fabrication of graphene–polymer
nanocomposites with higher-order three-dimensional architectures, Adv. Mater.,
2009, 21, 2180-2184.
[123] S. Stankovich, R. D. Piner, X. Chen, N. Wu, S. T. Nguyen, R. S. Ruoff,
Stable aqueous dispersions of graphitic nanoplatelets via the reduction of
exfoliated graphite oxide in the presence of poly(sodium 4–styrenesulfonate), J.
Mater. Chem., 2006, 16, 155.
133
[124] D. Li, M. B. Muller, S. Gilje, R. B. Kaner, G. G. Wallace, Processable
aqueous dispersions of graphene nanosheets, Nat. Nanotechnol., 2008, 3,
101-105.
[125] S. Stankovich, D. A. Dikin, R. D. Piner, K. A. Kohlhaas, A.
Kleinhammes, Y. Jia, Y. Wu, S. T. Nguyen, R. S. Ruoff, Synthesis of
graphene-based nanosheets via chemical reduction of exfoliated graphite oxide,
Carbon, 2007, 45, 1558-1565.
[126] X. Du, I. Skachko, A. Barker, E. Y. Andrei, Approaching ballistic
transport in suspended graphene, Nat. Nanotechnol., 2008, 3, 491-495.
[127] S. V. Morozov, K. S. Novoselov, M. I. Katsnelson, F. Schedin, D. C.
Elias, J. A. Jaszczak, A. K. Geim, Giant intrinsic carrier mobilities in graphene
and its bilayer, Phys. Rev. Lett., 2008, 100, 016602.
[128] Y. Lin, P. Avouris, Strong suppression of electrical noise in bilayer
graphene nanodevices, Nano Lett., 2008, 8, 2119-2125.
[129] K. I. Bolotin, K. J. Sikes, J. Hone, H. L. Stormer, P. Kim,
Temperature-dependent transport in suspended graphene, Phys. Rev. Lett., 2008,
101.
[130] A. Akturk, N. Goldsman, Electron transport and full-band
electron-phonon interactions in graphene, J. Appl. Phys., 2008, 103, 053702.
[131] F. V. Kusmartsev, W. M. Wu, M. P. Pierpoint, K. C. Yung,
Application of graphene within optoelectronic devices and transistors, in: P.
134
Misra (Ed.), Applied Spectroscopy and the Science of Nanomaterials, Springer
Singapore, 2015, pp. 191-221.
[132] C. Dimitrakopoulos, Y. Lin, A. Grill, D. B. Farmer, M. Freitag, Y. Sun,
S. Han, Z. Chen, K. A. Jenkins, Y. Zhu, Z. Liu, T. J. McArdle, J. A. Ott, R.
Wisnieff, P. Avouris, Wafer-scale epitaxial graphene growth on the Si-face of
hexagonal SiC (0001) for high frequency transistors, J. Vac. Sci. Technol. B,
2010, 28, 985-992.
[133] H. Feng, R. Cheng, X. Zhao, X. Duan, J. Li, A low-temperature
method to produce highly reduced graphene oxide, Nat. Commun., 2013, 4,
1539.
[134] I. W. Frank, D. M. Tanenbaum, A. M. van der Zande, P. L. McEuen,
Mechanical properties of suspended graphene sheets, J. Vac. Sci. Technol. B,
2007, 25, 2558-2561.
[135] C. Gómez-Navarro, M. Burghard, K. Kern, Elastic properties of
chemically derived single graphene sheets, Nano Lett., 2008, 8, 2045-2049.
[136] C. Dongyu, Y. Kamal, S. Mo, The mechanical properties and
morphology of a graphite oxide nanoplatelet/polyurethane composite,
Nanotechnology, 2009, 20, 085712.
[137] Y. Yang, C. Chen, Y. Wen, Q. Yang, M. Wang, Oxidized graphene
and graphene based polymer composites, New Carbon Mater., 2008, 23,
193-200.
135
[138] C. Yu, L. Shi, Z. Yao, D. Li, A. Majumdar, Thermal conductance and
thermopower of an individual single-wall carbon nanotube, Nano Lett., 2005, 5,
1842-1846.
[139] D. Nika, E. Pokatilov, A. Askerov, A. Balandin, Phonon thermal
conduction in graphene: Role of Umklapp and edge roughness scattering, Phys.
Rev. B, 2009, 79, 155413.
[140] Z. Guo, D. Zhang, X. Gong, Thermal conductivity of graphene
nanoribbons, Appl. Phys. Lett., 2009, 95, 163103.
[141] J. H. Seol, I. Jo, A. L. Moore, L. Lindsay, Z. H. Aitken, M. T. Pettes, X.
Li, Z. Yao, R. Huang, D. Broido, N. Mingo, R. S. Ruoff, L. Shi,
Two-dimensional phonon transport in supported graphene, Science, 2010, 328,
213-216.
[142] W. Cai, A. L. Moore, Y. Zhu, X. Li, S. Chen, L. Shi, R. S. Ruoff,
Thermal transport in suspended and supported monolayer graphene grown by
chemical vapor deposition, Nano Lett., 2010, 10, 1645-1651.
[143] T. Schwamb, B. R. Burg, N. C. Schirmer, D. Poulikakos, An electrical
method for the measurement of the thermal and electrical conductivity of
reduced graphene oxide nanostructures, Nanotechnology, 2009, 20, 405704.
[144] W. Yu, H. Xie, X. Wang, X. Wang, Significant thermal conductivity
enhancement for nanofluids containing graphene nanosheets, Phys. Lett. A,
2011, 375, 1323-1328.
136
[145] W. Chen, L. Yan, In situ self-assembly of mild chemical reduction
graphene for three-dimensional architectures, Nanoscale, 2011, 3, 3132-3137.
[146] I. Ivanov, A. Puretzky, G. Eres, H. Wang, Z. Pan, H. Cui, R. Jin, J.
Howe, D. B. Geohegan, Fast and highly anisotropic thermal transport through
vertically aligned carbon nanotube arrays, Appl. Phys. Lett., 2006, 89, 223110.
[147] V. A. Ermakov, A. V. Alaferdov, A. R. Vaz, A. V. Baranov, S. A.
Moshkalev, Nonlocal laser annealing to improve thermal contacts between
multi-layer graphene and metals, Nanotechnology, 2013, 24, 155301.
[148] Y. Zhang, Y. W. Tan, H. L. Stormer, P. Kim, Experimental
observation of the quantum Hall effect and Berry's phase in graphene, Nature,
2005, 438, 201-204.
[149] M. P. Levendorf, C. S. Ruiz-Vargas, S. Garg, J. Park, Transfer-free
batch fabrication of single layer graphene transistors, Nano Lett., 2009, 9,
4479-4483.
[150] Y. Lin, K. A. Jenkins, A. Valdes-Garcia, J. P. Small, D. B. Farmer, P.
Avouris, Operation of graphene transistors at Gigahertz frequencies, Nano Lett.,
2008, 9, 422-426.
[151] Y. Zhu, M. D. Stoller, W. Cai, A. Velamakanni, R. D. Piner, D. Chen,
R. S. Ruoff, Exfoliation of graphite oxide in propylene carbonate and thermal
reduction of the resulting graphene oxide platelets, ACS Nano, 2010, 4,
1227-1233.
137
[152] D. Tomanek, S. Berber, K. Umemoto, S. Saito, Hierarchical assembly
of nanostructured carbon foam, Mol. Cryst. Liq. Cryst. Sci. Technol., Sect. A,
2002, 386, 189-195.
[153] M. A. Worsley, P. J. Pauzauskie, S. O. Kucheyev, J. M. Zaug, A. V.
Hamza, J. H. Satcher, T. F. Baumann, Properties of single-walled carbon
nanotube-based aerogels as a function of nanotube loading, Acta Mater., 2009,
57, 5131-5136.
[154] M. A. Worsley, S. O. Kucheyev, J. D. Kuntz, T. Y. Olson, T. Y.-J. Han,
A. V. Hamza, J. H. Satcher, T. F. Baumann, Carbon scaffolds for stiff and
highly conductive monolithic oxide–carbon nanotube composites, Chem.
Mater., 2011, 23, 3054-3061.
[155] X. Zhang, J. Liu, B. Xu, Y. Su, Y. Luo, Ultralight conducting
polymer/carbon nanotube composite aerogels, Carbon, 2011, 49, 1884-1893.
[156] Y. Xu, K. Sheng, C. Li, G. Shi, Self-assembled graphene hydrogel via
a one-step hydrothermal process, ACS Nano, 2010, 4, 4324-4330.
[157] Z. Sui, X. Zhang, Y. Lei, Y. Luo, Easy and green synthesis of reduced
graphite oxide-based hydrogels, Carbon, 2011, 49, 4314-4321.
[158] U. M. Patil, J. S. Sohn, S. B. Kulkarni, S. C. Lee, H. G. Park, K. V.
Gurav, J. H. Kim, S. C. Jun, Enhanced supercapacitive performance of
chemically grown cobalt–nickel hydroxides on three-dimensional graphene
foam electrodes, ACS Appl. Mater. Interfaces, 2014, 6, 2450-2458.
138
[159] Z. Li, P. Liu, G. Yun, K. Shi, X. Lv, K. Li, J. Xing, B. Yang, 3D
(Three-dimensional) sandwich-structured of ZnO (zinc oxide)/rGO (reduced
graphene oxide)/ZnO for high performance supercapacitors, Energy, 2014, 69,
266-271.
[160] B. Tang, G. Hu, H. Gao, Z. Shi, Three-dimensional graphene network
assisted high performance dye sensitized solar cells, J. Power Sources, 2013,
234, 60-68.
[161] J. Wang, S. Liang, L. Ma, S. Ding, X. Yu, L. Zhou, Q. Wang, One-pot
synthesis of CdS-reduced graphene oxide 3D composites with enhanced
photocatalytic properties, Cryst. Eng. Comm., 2014, 16, 399-405.
[162] A. Lerf, H. He, M. Forster, J. Klinowski, Structure of graphite oxide
revisited, J. Phys. Chem. B, 1998, 102, 4477-4482.
[163] H. He, J. Klinowski, M. Forster, A. Lerf, A new structural model for
graphite oxide, Chem. Phys. Lett., 1998, 287, 53-56.
[164] G. Ruess, Über das graphitoxyhydroxyd (graphitoxyd), Monatsh.
Chem., 1947, 76, 381-417.
[165] T. Nakajima, Y. Matsuo, Formation process and structure of graphite
oxide, Carbon, 1994, 32, 469-475.
[166] T. Nakajima, A. Mabuchi, R. Hagiwara, A new structure model of
graphite oxide, Carbon, 1988, 26, 357-361.
139
[167] W. Cai, R. D. Piner, F. J. Stadermann, S. Park, M. A. Shaibat, Y. Ishii,
D. Yang, A. Velamakanni, S. J. An, M. Stoller, J. An, D. Chen, R. S. Ruoff,
Synthesis and solid-state NMR structural characterization of 13C-labeled
graphite oxide, Science, 2008, 321, 1815-1817.
[168] W. Gao, L. B. Alemany, L. Ci, P. M. Ajayan, New insights into the
structure and reduction of graphite oxide, Nature Chem., 2009, 1, 403-408.
[169] D. R. Dreyer, S. Park, C. W. Bielawski, R. S. Ruoff, The chemistry of
graphene oxide, Chem. Soc. Rev., 2010, 39, 228-240.
[170] M. Wissler, Graphite and carbon powders for electrochemical
applications, J. Power Sources, 2006, 156, 142-150.
[171] A. Lerf, H. He, T. Riedl, M. Forster, J. Klinowski, 13C and 1H MAS
NMR studies of graphite oxide and its chemically modified derivatives, Solid
State Ionics, 1997, 101-103, Part 2, 857-862.
[172] V. C. Tung, M. J. Allen, Y. Yang, R. B. Kaner, High-throughput
solution processing of large-scale graphene, Nat. Nanotechnol., 2009, 4, 25-29.
[173] Z. Fan, D. Z. Y. Tng, S. T. Nguyen, J. Feng, C. Lin, P. Xiao, L. Lu, H.
M. Duong, Morphology effects on electrical and thermal properties of
binderless graphene aerogels, Chem. Phys. Lett., 2013, 561-562, 92-96.
[174] S. T. Nguyen, H. T. Nguyen, A. Rinaldi, N. P. V. Nguyen, Z. Fan, H.
M. Duong, Morphology control and thermal stability of binderless-graphene
140
aerogels from graphite for energy storage applications, Colloids Surf., A, 2012,
414, 352-358.
[175] M. A. Worsley, T. Y. Olson, J. R. I. Lee, T. M. Willey, M. H. Nielsen,
S. K. Roberts, P. J. Pauzauskie, J. Biener, J. H. Satcher, T. F. Baumann, High
surface area, sp2-cross-linked three-dimensional graphene monoliths, J. Phys.
Chem. Lett., 2011, 2, 921-925.
[176] S. Zu, B. Han, Aqueous dispersion of graphene sheets stabilized by
pluronic copolymers: Formation of supramolecular hydrogel, J. Phys. Chem. C,
2009, 113, 13651-13657.
[177] Z. Tang, S. Shen, J. Zhuang, X. Wang, Noble-metal-promoted
three-dimensional macroassembly of single-layered graphene oxide, Angew.
Chem., Int. Ed. Engl., 2010, 49, 4603-4607.
[178] X. Jiang, Y. Ma, J. Li, Q. Fan, W. Huang, Self-assembly of reduced
graphene oxide into three-dimensional architecture by divalent ion linkage, J.
Phys. Chem. C, 2010, 114, 22462-22465.
[179] H. Vedala, D. C. Sorescu, G. P. Kotchey, A. Star, Chemical sensitivity
of graphene edges decorated with metal nanoparticles, Nano Lett., 2011, 11,
2342-2347.
[180] G. M. Scheuermann, L. Rumi, P. Steurer, W. Bannwarth, R. Mülhaupt,
Palladium nanoparticles on graphite oxide and its functionalized graphene
141
derivatives as highly active catalysts for the suzuki-miyaura coupling reaction,
J. Am. Chem. Soc., 2009, 131, 8262-8270.
[181] L. Zhang, G. Chen, M. N. Hedhili, H. Zhang, P. Wang,
Three-dimensional assemblies of graphene prepared by a novel chemical
reduction-induced self-assembly method, Nanoscale, 2012, 4, 7038-7045.
[182] S. Qin, X. Liu, R. Zhuo, X. Zhang, Microstructure-controllable
graphene oxide hydrogel film based on a pH-responsive graphene oxide
hydrogel, Macromol. Chem. Phys., 2012, 213, 2044-2051.
[183] M. T. Pettes, H. Ji, R. S. Ruoff, L. Shi, Thermal transport in
three-dimensional foam architectures of few-layer graphene and ultrathin
graphite, Nano Lett., 2012, 12, 2959-2964.
[184] K. Yang, J. Wang, B. Chen, Facile fabrication of stable monolayer and
few-layer graphene nanosheets as superior sorbents for persistent aromatic
pollutant management in water, J. Mater. Chem. A, 2014, 2, 18219-18224.
[185] M. Mecklenburg, A. Schuchardt, Y. K. Mishra, S. Kaps, R. Adelung,
A. Lotnyk, L. Kienle, K. Schulte, Aerographite: Ultra lightweight, flexible
nanowall, carbon microtube material with outstanding mechanical performance,
Adv. Mater., 2012, 24, 3486-3490.
[186] S. Ye, J. Feng, P. Wu, Highly elastic graphene oxide–epoxy composite
aerogels via simple freeze-drying and subsequent routine curing, J. Mater.
Chem. A, 2013, 1, 3495-3502.
142
[187] Y. He, Y. Liu, T. Wu, J. Ma, X. Wang, Q. Gong, W. Kong, F. Xing, Y.
Liu, J. Gao, An environmentally friendly method for the fabrication of reduced
graphene oxide foam with a super oil absorption capacity, J. Hazard. Mater.,
2013, 260, 796-805.
[188] M. A. Worsley, P. J. Pauzauskie, T. Y. Olson, J. Biener, J. H. Satcher,
T. F. Baumann, Synthesis of graphene aerogel with high electrical conductivity,
J. Am. Chem. Soc., 2010, 132, 14067-14069.
[189] Z. M. Marković, B. M. Babić, M. D. Dramićanin, I. D. Holclajtner
Antunović, V. B. Pavlović, D. B. Peruško, B. M. Todorović Marković,
Preparation of highly conductive carbon cryogel based on pristine graphene,
Synth. Met., 2012, 162, 743-747.
[190] S. Pei, H. Cheng, The reduction of graphene oxide, Carbon, 2012, 50,
3210-3228.
[191] H. Huang, P. Chen, X. Zhang, Y. Lu, W. Zhan, Edge-to-edge
assembled graphene oxide aerogels with outstanding mechanical performance
and superhigh chemical activity, Small, 2013, 9, 1397-1404.
[192] Y. Zhong, M. Zhou, F. Huang, T. Lin, D. Wan, Effect of graphene
aerogel on thermal behavior of phase change materials for thermal management,
Sol. Energy Mater. Sol. Cells, 2013, 113, 195-200.
143
[193] A. J. McNamara, Y. Joshi, Z. M. Zhang, Characterization of
nanostructured thermal interface materials–A review, Int. J. Therm. Sci., 2012,
62, 2-11.
[194] K. M. Shahil, A. A. Balandin, Graphene-multilayer graphene
nanocomposites as highly efficient thermal interface materials, Nano Lett.,
2012, 12, 861-867.
[195] B. Debelak, K. Lafdi, Use of exfoliated graphite filler to enhance
polymer physical properties, Carbon, 2007, 45, 1727-1734.
[196] J. Xu, T. S. Fisher, Enhancement of thermal interface materials with
carbon nanotube arrays, Int. J. Heat Mass Transfer, 2006, 49, 1658-1666.
[197] P. Bonnet, D. Sireude, B. Garnier, O. Chauvet, Thermal properties and
percolation in carbon nanotube–polymer composites, Appl. Phys. Lett., 2007,
91, 201910.
[198] M. T. Hung, O. Choi, Y. S. Ju, H. T. Hahn, Heat conduction in
graphite–nanoplatelet–reinforced polymer nanocomposites, Appl. Phys. Lett.,
2006, 89, 023117.
[199] X. J. Hu, A. A. Padilla, J. Xu, T. S. Fisher, K. E. Goodson, 3-omega
measurements of vertically oriented carbon nanotubes on silicon, J. Heat
Transfer, 2006, 128, 1109-1113.
144
[200] X. J. Hu, M. A. Panzer, K. E. Goodson, Infrared microscopy thermal
characterization of opposing carbon nanotube arrays, J. Heat Transfer, 2007,
129, 91-93.
[201] A. McNamara, V. Sahu, Y. Joshi, Z. Zhang, Infrared imaging
microscope as an effective tool for measuring thermal resistance of emerging
interface materials, ASME, Amer Soc Mechanical Engineers, Chicago, 2011,
pp. 189-194.
[202] Y. Sun, Q. Wu, G. Shi, Supercapacitors based on self-assembled
graphene organogel, Phys. Chem. Chem. Phys., 2011, 13, 17249-17254.
[203] L. Zhang, G. Shi, Preparation of highly conductive graphene hydrogels
for fabricating supercapacitors with high rate capability, J. Phys. Chem. C, 2011,
115, 17206-17212.
[204] R. Li, C. Chen, J. Li, L. Xu, G. Xiao, D. Yan, A facile approach to
superhydrophobic and superoleophilic graphene/polymer aerogels, J. Mater.
Chem. A, 2014, 2, 3057.
[205] X. Gui, Z. Zeng, Y. Zhu, H. Li, Z. Lin, Q. Gan, R. Xiang, A. Cao, Z.
Tang, Three-dimensional carbon nanotube sponge–array architectures with
high energy dissipation, Adv. Mater., 2014, 26, 1248-1253.
[206] W. Yang, K. R. Ratinac, S. P. Ringer, P. Thordarson, J. J. Gooding, F.
Braet, Carbon nanomaterials in biosensors: Should you use nanotubes or
graphene?, Angew. Chem., Int. Ed. Engl., 2010, 49, 2114-2138.
145
[207] S. H. Lee, D. H. Lee, W. J. Lee, S. O. Kim, Tailored assembly of
carbon nanotubes and graphene, Adv. Funct. Mater., 2011, 21, 1338-1354.
[208] K. H. Kim, Y. Oh, M. F. Islam, Graphene coating makes carbon
nanotube aerogels superelastic and resistant to fatigue, Nat. Nanotechnol., 2012,
7, 562-566.
[209] Y. Wang, Y. Wu, Y. Huang, F. Zhang, X. Yang, Y. Ma, Y. Chen,
Preventing graphene sheets from restacking for high-capacitance performance,
J. Phys. Chem. C, 2011, 115, 23192-23197.
[210] E. Yoo, J. Kim, E. Hosono, H. Zhou, T. Kudo, I. Honma, Large
reversible Li storage of graphene nanosheet families for use in rechargeable
lithium ion batteries, Nano Lett., 2008, 8, 2277-2282.
[211] V. C. Tung, L. Chen, M. J. Allen, J. K. Wassei, K. Nelson, R. B. Kaner,
Y. Yang, Low-temperature solution processing of graphene-carbon nanotube
hybrid materials for high-performance transparent conductors, Nano Lett., 2009,
9, 1949-1955.
[212] Z. Sui, Q. Meng, X. Zheng, R. Ma, C. Bing, Green synthesis of carbon
nanotube–graphene hybrid aerogels and their use as versatile agents for water
purification, J. Mater. Chem., 2012, 22, 8767.
[213] X. Dong, J. Chen, Y. Ma, J. Wang, M. B. Chan-Park, X. Liu, L. Wang,
W. Huang, P. Chen, Superhydrophobic and superoleophilic hybrid foam of
146
graphene and carbon nanotube for selective removal of oils or organic solvents
from the surface of water, Chem. Commun., 2012, 48, 10660-10662.
[214] C. Bao, Y. Guo, L. Song, Y. Hu, Poly(vinyl alcohol) nanocomposites
based on graphene and graphite oxide: A comparative investigation of property
and mechanism, J. Mater. Chem., 2011, 21, 13942-13950.
[215] X. Yang, L. Li, S. Shang, X. Tao, Synthesis and characterization of
layer-aligned poly(vinyl alcohol)/graphene nanocomposites, Polymer, 2010, 51,
3431-3435.
[216] J. R. Potts, S. H. Lee, T. M. Alam, J. An, M. D. Stoller, R. D. Piner, R.
S. Ruoff, Thermomechanical properties of chemically modified
graphene/poly(methyl methacrylate) composites made by in situ
polymerization, Carbon, 2011, 49, 2615-2623.
[217] W. Li, X. Tang, H. Zhang, Z. Jiang, Z. Yu, X. Du, Y. Mai,
Simultaneous surface functionalization and reduction of graphene oxide with
octadecylamine for electrically conductive polystyrene composites, Carbon,
2011, 49, 4724-4730.
[218] Y. S. Yun, Y. H. Bae, D. H. Kim, J. Y. Lee, I. Chin, H. Jin, Reinforcing
effects of adding alkylated graphene oxide to polypropylene, Carbon, 2011, 49,
3553-3559.
147
[219] C. Teng, C. M. Ma, C. Lu, S. Yang, S. Lee, M. Hsiao, M. Yen, K.
Chiou, T. Lee, Thermal conductivity and structure of non-covalent
functionalized graphene/epoxy composites, Carbon, 2011, 49, 5107-5116.
[220] J. Liang, Y. Xu, Y. Huang, L. Zhang, Y. Wang, Y. Ma, F. Li, T. Guo, Y.
Chen, Infrared-triggered actuators from graphene-based nanocomposites, J.
Phys. Chem. C, 2009, 113, 9921-9927.
[221] J. Yan, T. Wei, B. Shao, Z. Fan, W. Qian, M. Zhang, F. Wei,
Preparation of a graphene nanosheet/polyaniline composite with high specific
capacitance, Carbon, 2010, 48, 487-493.
[222] M. T. Pettes, I. Jo, Z. Yao, L. Shi, Influence of polymeric residue on
the thermal conductivity of suspended bilayer graphene, Nano Lett., 2011, 11,
1195-1200.
[223] H. Wang, J. T. Robinson, X. Li, H. Dai, Solvothermal reduction of
chemically exfoliated graphene sheets, J. Am. Chem. Soc., 2009, 131,
9910-9911.
[224] X. Zeng, J. Yang, W. Yuan, Preparation of a poly(methyl
methacrylate)–reduced graphene oxide composite with enhanced properties by
a solution blending method, Eur. Polym. J., 2012, 48, 1674-1682.
[225] Y. Li, J. Zhu, S. Wei, J. Ryu, L. Sun, Z. Guo,
Poly(propylene)/graphene nanoplatelet nanocomposites: Melt rheological
148
behavior and thermal, electrical, and electronic properties, Macromol. Chem.
Phys., 2011, 212, 1951-1959.
[226] J. Yu, P. Jiang, C. Wu, L. Wang, X. Wu, Graphene nanocomposites
based on poly(vinylidene fluoride): Structure and properties, Polym. Compos.,
2011, 32, 1483-1491.
[227] T. Ramanathan, A. A. Abdala, S. Stankovich, D. A. Dikin, M.
Herrera-Alonso, R. D. Piner, D. H. Adamson, H. C. Schniepp, X. Chen, R. S.
Ruoff, S. T. Nguyen, I. A. Aksay, R. K. Prud'Homme, L. C. Brinson,
Functionalized graphene sheets for polymer nanocomposites, Nat.
Nanotechnol., 2008, 3, 327-331.
[228] H. Kim, Y. Miura, C. W. Macosko, Graphene/polyurethane
nanocomposites for improved gas barrier and electrical conductivity, Chem.
Mater., 2010, 22, 3441-3450.
[229] H. Zhang, W. Zheng, Q. Yan, Y. Yang, J. Wang, Z. Lu, G. Ji, Z. Yu,
Electrically conductive polyethylene terephthalate/graphene nanocomposites
prepared by melt compounding, Polymer, 2010, 51, 1191-1196.
[230] H. Kim, C. W. Macosko, Processing-property relationships of
polycarbonate/graphene composites, Polymer, 2009, 50, 3797-3809.
[231] M. A. Aldosari, A. A. Othman, E. H. Alsharaeh, Synthesis and
characterization of the in situ bulk polymerization of PMMA containing
graphene sheets using microwave irradiation, Molecules, 2013, 18, 3152-3167.
149
[232] J. Wang, Z. Shi, Y. Ge, Y. Wang, J. Fan, J. Yin, Solvent exfoliated
graphene for reinforcement of PMMA composites prepared by in situ
polymerization, Mater. Chem. Phys., 2012, 136, 43-50.
[233] T. Kuila, S. Bose, P. Khanra, N. H. Kim, K. Y. Rhee, J. H. Lee,
Characterization and properties of in situ emulsion polymerized poly(methyl
methacrylate)/graphene nanocomposites, Compos. Part. A: Appl. S, 2011, 42,
1856-1861.
[234] Y. Guo, C. Bao, L. Song, B. Yuan, Y. Hu, In situ polymerization of
graphene, graphite oxide, and functionalized graphite oxide into epoxy resin
and comparison study of on-the-flame behavior, Ind. Eng. Chem. Res., 2011, 50,
7772-7783.
[235] Y. R. Lee, A. V. Raghu, H. M. Jeong, B. K. Kim, Properties of
waterborne polyurethane/functionalized graphene sheet nanocomposites
prepared by an in situ method, Macromol. Chem. Phys., 2009, 210, 1247-1254.
[236] S. Stankovich, D. A. Dikin, G. H. Dommett, K. M. Kohlhaas, E. J.
Zimney, E. A. Stach, R. D. Piner, S. T. Nguyen, R. S. Ruoff, Graphene-based
composite materials, Nature, 2006, 442, 282-286.
[237] V. H. Pham, T. V. Cuong, T. T. Dang, S. H. Hur, B. Kong, E. J. Kim, E.
W. Shin, J. S. Chung, Superior conductive polystyrene–chemically converted
graphene nanocomposite, J. Mater. Chem., 2011, 21, 11312-11316.
150
[238] J. Liang, Y. Huang, L. Zhang, Y. Wang, Y. Ma, T. Guo, Y. Chen,
Molecular-level dispersion of graphene into poly(vinyl alcohol) and effective
reinforcement of their nanocomposites, Adv. Funct. Mater., 2009, 19,
2297-2302.
[239] J. Liang, Y. Wang, Y. Huang, Y. Ma, Z. Liu, J. Cai, C. Zhang, H. Gao,
Y. Chen, Electromagnetic interference shielding of graphene/epoxy composites,
Carbon, 2009, 47, 922-925.
[240] H. Pang, T. Chen, G. Zhang, B. Zeng, Z. Li, An electrically conducting
polymer/graphene composite with a very low percolation threshold, Mater.
Lett., 2010, 64, 2226-2229.
[241] T. D. Dao, G. Erdenedelger, H. M. Jeong, Water-dispersible graphene
designed as a Pickering stabilizer for the suspension polymerization of
poly(methyl methacrylate)/graphene core–shell microsphere exhibiting
ultra-low percolation threshold of electrical conductivity, Polymer, 2014, 55,
4709-4719.
[242] Z. Chen, W. Ren, L. Gao, B. Liu, S. Pei, H. Cheng, Three-dimensional
flexible and conductive interconnected graphene networks grown by chemical
vapour deposition, Nat. Mater., 2011, 10, 424-428.
[243] J. Wang, S. Yang, Y. Huang, H. Tien, W. Chin, C. M. Ma, Preparation
and properties of graphene oxide/polyimide composite films with low dielectric
151
constant and ultrahigh strength via in situ polymerization, J. Mater. Chem.,
2011, 21, 13569-13575.
[244] J. Wang, X. Wang, C. Xu, M. Zhang, X. Shang, Preparation of
graphene/poly(vinyl alcohol) nanocomposites with enhanced mechanical
properties and water resistance, Polym. Int., 2011, 60, 816-822.
[245] H. Zhang, Q. Yan, W. Zheng, Z. He, Z. Yu, Tough graphene−polymer
microcellular foams for electromagnetic interference shielding, ACS Appl.
Mater. Interfaces, 2011, 3, 918-924.
[246] C. K. Leong, D. D. L. Chung, Carbon black dispersions as thermal
pastes that surpass solder in providing high thermal contact conductance,
Carbon, 2003, 41, 2459-2469.
[247] A. A. Balandin, Thermal properties of graphene and nanostructured
carbon materials, Nat. Mater., 2011, 10, 569-581.
[248] F. Yavari, H. R. Fard, K. Pashayi, M. A. Rafiee, A. Zamiri, Z. Yu, R.
Ozisik, T. Borca-Tasciuc, N. Koratkar, Enhanced thermal conductivity in a
nanostructured phase change composite due to low concentration graphene
additives, J. Phys. Chem. C, 2011, 115, 8753-8758.
[249] A. Yu, P. Ramesh, M. E. Itkis, E. Bekyarova, R. C. Haddon,
pollackGraphite nanoplatelet–epoxy composite thermal interface materials, J.
Phys. Chem. C, 2007, 111, 7565-7569.
[250] G. Pollack, Kapitza resistance, Rev. Mod. Phys., 1969, 41, 48-81.
152
[251] C. W. Nan, R. Birringer, D. R. Clarke, H. Gleiter, Effective thermal
conductivity of particulate composites with interfacial thermal resistance, J.
Appl. Phys., 1997, 81, 6692-6699.
[252] J. Yan, T. Wei, Z. Fan, W. Qian, M. Zhang, X. Shen, F. Wei,
Preparation of graphene nanosheet/carbon nanotube/polyaniline composite as
electrode material for supercapacitors, J. Power Sources, 2010, 195,
3041-3045.
[253] S. Stankovich, R. D. Piner, S. T. Nguyen, R. S. Ruoff, Synthesis and
exfoliation of isocyanate-treated graphene oxide nanoplatelets, Carbon, 2006,
44, 3342-3347.
[254] L. Q. Xu, Y. L. Liu, K. G. Neoh, E. T. Kang, G. D. Fu, Reduction of
graphene oxide by aniline with its concomitant oxidative polymerization,
Macromol. Rapid Commun., 2011, 32, 684-688.
[255] J. Als-Nielsen, D. McMorrow, X-rays and their interaction with matter,
Elements of Modern X-ray Physics, John Wiley & Sons, Inc., 2011, pp. 1-28.
[256] D. Newbury, D. Joy, P. Echlin, C. Fiori, J. Goldstein, Electron
channeling contrast in the SEM, Advanced Scanning Electron Microscopy and
X-Ray Microanalysis, Springer US, 1986, pp. 87-145.
[257] G. Fagerlund, Determination of specific surface by the BET method,
Matériaux et Construction, 1973, 6, 239-245.
153
[258] R. Radhakrishnan, A review of: “Thermoelectric handbook, macro to
nano, D.M. Rowe (editor)”, Mater. Manuf. Processes, 2008, 23, 626-627.
[259] A. W. Coats, J. P. Redfern, Thermogravimetric analysis: A review,
Analyst, 1963, 88, 906-924.
[260] I. H. Buckle, Progress in micro-indentation hardness testing, Int. Mater.
Rev., 1959, 4, 49-100.
[261] J. Zhang, H. Yang, G. Shen, P. Cheng, J. Zhang, S. Guo, Reduction of
graphene oxide via L–ascorbic acid, Chem. Commun., 2010, 46, 1112-1114.
[262] P. Lian, X. Zhu, S. Liang, Z. Li, W. Yang, H. Wang, Large reversible
capacity of high quality graphene sheets as an anode material for lithium-ion
batteries, Electrochim. Acta, 2010, 55, 3909-3914.
[263] K. Muthoosamy, R. G. Bai, I. B. Abubakar, S. M. Sudheer, H. N. Lim,
H.-S. Loh, N. M. Huang, C. H. Chia, S. Manickam, Exceedingly biocompatible
and thin-layered reduced graphene oxide nanosheets using an eco-friendly
mushroom extract strategy, Int. J. Nanomedicine, 2015, 10, 1505-1519.
[264] W. Hayes, P. Joseph, M. Mughal, P. Papakonstantinou, Production of
reduced graphene oxide via hydrothermal reduction in an aqueous sulphuric
acid suspension and its electrochemical behaviour, J. Solid State Electrochem.,
2015, 19, 361-380.
[265] P. Wang, T. Hsieh, C. Chiang, M. Shen, Synergetic effects of
mechanical properties on graphene nanoplatelet and multiwalled carbon
154
nanotube hybrids reinforced epoxy/carbon fiber composites, J. Nanomater.,
2015, Article ID: 838032, In press.
[266] K. Sheng, Y. Xu, C. Li, G. Shi, High-performance self-assembled
graphene hydrogels prepared by chemical reduction of graphene oxide, New
Carbon Mater., 2011, 26, 9-15.
[267] G. Chester, A. Thellung, The law of Wiedemann and Franz, Proc. Phys.
Soc., 1961, 77, 1005.
[268] E. Pop, V. Varshney, A. K. Roy, Thermal properties of graphene:
Fundamentals and applications, MRS Bull., 2012, 37, 1273-1281.
[269] R. Singh, H. S. Kasana, Computational aspects of effective thermal
conductivity of highly porous metal foams, Appl. Therm. Eng., 2004, 24,
1841-1849.
[270] S. N. Schiffres, K. H. Kim, L. Hu, A. J. H. McGaughey, M. F. Islam, J.
A. Malen, Gas diffusion, energy transport, and thermal accommodation in
single-walled carbon nanotube aerogels, Adv. Funct. Mater., 2012, 22,
5251-5258.
[271] J. Y. Kim, J. Lee, J. C. Grossman, Thermal transport in functionalized
graphene, ACS Nano, 2012, 6, 9050-9057.
[272] C. Gomez-Navarro, J. C. Meyer, R. S. Sundaram, A. Chuvilin, S.
Kurasch, M. Burghard, K. Kern, U. Kaiser, Atomic structure of reduced
graphene oxide, Nano Lett., 2010, 10, 1144-1148.
155
[273] S. Ghosh, W. Bao, D. L. Nika, S. Subrina, E. P. Pokatilov, C. N. Lau, A.
A. Balandin, Dimensional crossover of thermal transport in few-layer graphene,
Nat. Mater., 2010, 10, 555-558.
[274] W. D. Kingery, Thermal Conductivity: XIV, conductivity of
multi-component systems, J. Am. Ceram. Soc., 1959, 42, 617-627.
[275] H. D. Wang, J. H. Lin, Z. Y. Guo, X. Zhang, R. F. Zhang, F. Wei, T. Y.
Li, Thermal transport across the interface between a suspended single-walled
carbon nanotube and air, Nanosc. Microsc. Therm., 2013, 17, 349-365.
[276] Y. Zhang, L. Liu, N. Xi, Y. Wang, Z. Dong, U. C. Wejinya,
Dielectrophoretic assembly and atomic force microscopy modification of
reduced graphene oxide, J. Appl. Phys., 2011, 110, 114515.
[277] K. M. F. Shahil, A. A. Balandin, Thermal properties of graphene and
multilayer graphene: Applications in thermal interface materials, Solid State
Commun., 2012, 152, 1331-1340.
[278] T. Ramanathan, S. Stankovich, D. A. Dikin, H. Liu, H. Shen, S. T.
Nguyen, L. C. Brinson, Graphitic nanofillers in PMMA nanocomposites–An
investigation of particle size and dispersion and their influence on
nanocomposite properties, J. Polym. Sci., Part B: Polym. Phys., 2007, 45,
2097-2112.
156
[279] H. J. Salavagione, G. Martínez, M. A. Gómez, Synthesis of poly(vinyl
alcohol)/reduced graphite oxide nanocomposites with improved thermal and
electrical properties, J. Mater. Chem., 2009, 19, 5027.
[280] W. Zheng, S. Wong, Electrical conductivity and dielectric properties
of PMMA/expanded graphite composites, Compos. Sci. Technol., 2003, 63,
225-235.
[281] H. Chakraborty, A. Sinha, N. Mukherjee, P. P. Chattopadhyay,
Exfoliated graphite reinforced PMMA composite: A study on nanoindentation
and scratch behavior, J. Nanotechnol., 2012, 2012, 1-5.
[282] F. Bahrani, A. Safari, V. Mahroo;, K. Ghasem, Comparison of
hardness and surface roughness of two denture bases polymerized by different
methods, World J. Dent., 2012, 3, 171-175.
[283] S. R. Tatro, L. M. Clayton, P. A. O'Rourke Muisener, A. M. Rao, J. P.
Harmon, Probing multi-walled nanotube/poly(methyl methacrylate)
composites with ionizing radiation, Polymer, 2004, 45, 1971-1979.
[284] G. H. Chen, D. J. Wu, W. G. Weng, W. L. Yan, Preparation of
polymer/graphite conducting nanocomposite by intercalation polymerization, J.
Appl. Polym. Sci., 2001, 82, 2506-2513.
[285] Y. Feng, H. Zheng, Z. Zhu, F. Zu, The microstructure and electrical
conductivity of aluminum alloy foams, Mater. Chem. Phys., 2002, 78, 196-201.
157
[286] X. Y. Yuan, Improved properties of chemically modified
graphene/poly(methyl methacrylate) nanocomposites via a facile in situ bulk
polymerization, Express Polym. Lett., 2012, 6, 847-858.
[287] R. Pylkki, P. Moyer, P. West, Scanning near-field optical microscopy
and scanning thermal microscopy, Jpn. J. Appl. Phys., 1994, 33, 3785-3790.
[288] J. Brandrup, E. H. Immergut, E. A. Grulke, Polymer data handbook,
John Wiley & Sons, Inc., 1999, New York V/88.
[289] H. M. Duong, D. V. Papavassiliou, K. J. Mullen, S. Maruyama,
Computational modeling of the thermal conductivity of single-walled carbon
nanotube–polymer composites, Nanotechnology, 2008, 19, 065702.
[290] S. Heo, S. Y. Cho, D. H. Kim, Y. Choi, H. H. Park, H. Jin, Improved
thermal properties of graphene oxide–incorporated poly(methyl methacrylate)
microspheres, J. Nanosci. Nanotechnol., 2012, 12, 5990-5994.
[291] D. Konatham, D. V. Papavassiliou, A. Striolo, Thermal boundary
resistance at the graphene–graphene interface estimated by molecular dynamics
simulations, Chem. Phys. Lett., 2012, 527, 47-50.
[292] Z. Fan, A. Marconnet, S. T. Nguyen, C. Y. H. Lim, H. M. Duong,
Effects of heat treatment on the thermal properties of highly nanoporous
graphene aerogels using the infrared microscopy technique, Int. J. Heat Mass
Transfer, 2014, 76, 122-127.
158
[293] B. Das, K. Eswar Prasad, U. Ramamurty, C. N. Rao, Nano-indentation
studies on polymer matrix composites reinforced by few-layer graphene,
Nanotechnology, 2009, 20, 125705.
[294] M. M. Shokrieh, M. R. Hosseinkhani, M. R. Naimi-Jamal, H. Tourani,
Nanoindentation and nanoscratch investigations on graphene-based
nanocomposites, Polym. Test., 2013, 32, 45-51.
[295] D. Blond, V. Barron, M. Ruether, K. P. Ryan, V. Nicolosi, W. J. Blau,
J. N. Coleman, Enhancement of modulus, strength, and toughness in
poly(methyl methacrylate)-based composites by the Incorporation of
poly(methyl methacrylate)–functionalized Nanotubes, Adv. Funct. Mater.,
2006, 16, 1608-1614.
[296] H. Hu, L. Onyebueke, A. Abatan, Characterizing and modeling
mechanical properties of nanocomposites–Review and evaluation, JMMCE,
2010, 9, 275-319.
[297] J. A. King, D. R. Klimek, I. Miskioglu, G. M. Odegard, Mechanical
properties of graphene nanoplatelet/epoxy composites, J. Appl. Polym. Sci.,
2013, 128, 4217-4223.
[298] T. Forati, M. Atai, A. M. Rashidi, M. Imani, A. Behnamghader,
Physical and mechanical properties of graphene oxide/polyethersulfone
nanocomposites, Polym. Adv. Technol., 2014, 25, 322-328.
159
[299] M. Moazzami Gudarzi, F. Sharif, Enhancement of dispersion and
bonding of graphene–polymer through wet transfer of functionalized graphene
oxide, Express Polym. Lett., 2012, 6, 1017-1031.
[300] B. Phillip, J. K. Abraham, A. Chandrasekhar, V. K. Varadan, Carbon
nanotube/PMMA composite thin films for gas-sensing applications, Smart
Mater. Struct., 2003, 12, 935-939.
[301] A. Pandey, PMMA–CNT matrices for vacuum electronic, biosensing
and energy applications, Department of Physics, Michigan Technological
University, 2010.
160
LIST OF PUBLICATIONS
1. BOOK CHAPTER
(1) Hai M. Duong, Zeng Fan, Son T. Nguyen, Carbon Nanotube/Graphene
Aerogels, The Carbon Nanomaterials Sourcebook, Taylor & Francis (CRC
Press), 2015.
2. JOURNAL PAPERS
(1) Zeng Fan, Feng Gong, Son T. Nguyen, Hai M. Duong, Advanced
multifunctional graphene aerogel–poly (methyl methacrylate) composites:
Experiments and modeling, Carbon, 2015, 81, 396-404.
(2) Zeng Fan, Amy Marconnet, Son T. Nguyen, Christina Y. H. Lim, Hai M.
Duong, Effects of heat treatment on the thermal properties of highly
nanoporous graphene aerogels using the infrared microscopy technique,
International Journal of Heat and Mass Transfer, 2014, 76, 122-127.
(3) Zeng Fan, Daniel Z. Y. Tng, Clarisse X. T. Lim, Peng Liu, Son T. Nguyen,
Pengfei Xiao, Amy Marconnet, Christina Y. H. Lim, Hai M. Duong,
Thermal and electrical properties of graphene/carbon nanotube aerogels,
Colloids and Surfaces A: Physicochemical and Engineering Aspects, 2014,
445, 48-53.
(4) Zeng Fan, Daniel Z. Y. Tng, Son T. Nguyen, Chunfu Lin, Pengfei Xiao, Li
Lu, Hai M. Duong, Morphology effects on electrical and thermal properties
of binderless graphene aerogels, Chemical Physics Letters, 2013, 561–562,
92-96.
161
(5) Son T. Nguyen, Hoa T. Nguyen, Ali Rinaldi, Nam P. V. Nguyen, Zeng Fan,
Hai M. Duong, Morphology control and thermal stability of
binderless-graphene aerogels from graphite for energy storage applications,
Colloids and Surfaces A: Physicochemical and Engineering Aspects, 2012,
414, 352-358.
(6) Jingduo Feng, Son T. Nguyen, Zeng Fan, Hai M. Duong, Advanced
fabrication and oil absorption properties of superhydrophobic recycled
cellulose aerogels, Chemical Engineering Journal, 2015, revised.
3. CONFERENCE PROCEEDINGS
(1) Zeng Fan, Peng Liu, Thang Q. Tran, Hai M. Duong, Advanced fabrication
of carbon nanotube aerogels by chemical vapor deposition. NT15, 29th
June-3rd
July 2015, Nagoya, Japan.
(2) Thang Q. Tran, Peng Liu, Zeng Fan, Nigel H.H. Ngern, Hai M. Duong,
Advanced multifunctional properties of aligned carbon nanotube–epoxy
composites from carbon nanotube aerogel method, 2015 MRS Spring
Meeting & Exhibit, 6th
-10th
April 2015, San Francisco, California, USA.
(3) Peng Liu, Thang Q. Tran, Zeng Fan, Nigel H. H. Ngern, Hai M. Duong,
Advanced multifunctional properties of aligned carbon nanotube–epoxy
composites from carbon nanotube aerogel method, APS March Meeting
2015, 2nd
-6th
March 2015, San Antonio, Texas, USA.
(4) Peng Liu, Zeng Fan, Abhik Damani, Thang Q. Tran, Hanlin Cheng, Son T.
Nguyen, Vincent Tan, Hai M. Duong, Synthesis of carbon
nanotube/graphene aerogel hybrid material by chemical vapor deposition
162
for energy storage devices, International Meeting on the Chemistry of
Graphene and Carbon Nanotubes, 30th
March-3rd
April 2014, Riva del
Garda, Italy.
(5) Zeng Fan, Daniel Z. Y. Tng, Clarisse X. T. Lim, Son T. Nguyen, Amy
Marconnet, Christina Y. H. Lim, Hai M. Duong, Electrical and thermal
properties of morphology–controlled carbon nanotube/graphene aerogels,
International Meeting on the Chemistry of Graphene and Carbon Nanotubes,
30th
March-3rd
April 2014, Riva del Garda, Italy.
(6) Peng Liu, Zeng Fan, Thang Q. Tran, Hanlin Cheng, Son T. Nguyen, Hai M.
Duong, Vertically aligned carbon nanotubes/graphene aerogels for energy
storage devices, 22nd
International Conference on Processing and
Fabrication of Advanced Materials (PFAM XXII), 18th
-20th
December
2013, Kent Ridge Guild House, Singapore.
(7) Hanlin Cheng, Stephanie C. P. Neo, Lilian Medina, Son T. Nguyen, Zeng
Fan, Hai M. Duong, Carbon/MWNTs and PEDOT/MWNTs aerogel
composites for supercapacitor application, International Conference on
Materials for Advanced Technologies (ICMAT). 30th
June-5th
July 2013,
Singapore.
(8) Zeng Fan, Daniel Z. Y. Tng, Clarisse X. T. Lim, Chunfu Lin, Son T.
Nguyen, Amy Marconnet, Hai M. Duong, Electrical and thermal properties
of morphology–controlled graphene/carbon nanotube aerogels for energy
storage devices, NT13 : Fourteen International Conference on the Science
and Applications of Nanotubes, 24th
-28th
June 2013, Espoo, Finland.
163
(9) Zeng Fan, Daniel Z. Y. Tng, Son T. Nguyen, Jingduo Feng, Hai M. Duong.
Morphology effects on electrical properties of advanced graphene aerogels,
2013 MRS Spring Meeting & Exhibit, 1st-5
th April 2013, San Francisco,
California, USA.
(10) Shao Kai Ng, Jia Cheng Oh, Janet P. W. Wong, Son T. Nguyen, Zeng Fan,
Hai M. Duong, Green recycled cellulose aerogels from waste paper for oil
spill cleaning, 2013 MRS Spring Meeting & Exhibit, 1st-5
th April 2013, San
Francisco, California, USA.
(11) Zeng Fan, Son T. Nguyen, Daniel Z. Y. Tng, Clarisse X. T. Lim, Jingduo
Feng, Stephanie C. P. Neo, Hai M. Duong, Simple but effective method of
morphology control of graphene aerogels for energy applications, Annual
International Conference on Chemistry, Chemical Engineering and
Chemical Process (CCECP 2013), 25th
-26th
February 2013, Singapore.
