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Chapter 1
Introduction and Literature Review
1.1 Background
Nanomaterials belong to a new branch of technology which has surpassed all
conventional barriers and boundaries that differentiated science from engineering to
bring about the confluence of the two. It has, perhaps for the first time since centuries,
created a new platform for scientists and engineers all over the world to come and
work together through strong linkages, networking and collaborations. Hence path
breaking and highly innovative technologies and concepts have emerged during the
last decade.
The nanoscale and associated nanoscience and technology have afforded
unique opportunities to create revolutionary material combinations. Naturally-
occurring biological systems have taken advantage of the properties of the
interactions of organic materials at the nanoscale and at present, work is in progress
for exploiting combinations of inorganic materials. These new generation materials
circumvent many of the classic material performance trade-offs by accessing some
desirable properties and exploiting unique synergisms between materials. This has
been possible at length-scale of morphology and the fundamental physics associated
with properties coincide on the nanoscale. The combination of fundamental
understanding of materials and the realization of fabrication and processing
techniques provide simultaneous structural control on the nano, micro and macro-
length scales. It is perhaps the heart of nanoengineered materials. Availability of such
material of modern technologies are rapidly increasing, impacting many diverse areas
of commercial applications and military arena.
Epoxy resin is a thermoset polymer that contains two or more epoxide groups.
It is one of the most commonly used thermosetting micro molecular synthetic
material. Epoxy resins have properties of excellent mechanical strength, dielectric
property and chemical stability [1]. It has additional advantages such as low
concentration percentage, low cost, easy mouldability etc. It is widely used in many
applications as a composite material for insulation, anti corrosion coating,
cementation between metals and non-metals and structural applications. In fact, epoxy
based composites have become one of the most important materials for engineering
applications.
In the past, research on modification of epoxy systems was limited to rubber.
However, with increase in research activities on epoxy resin, new methods of epoxy
resin modifications such as liquid crystal and nanoparticles were introduced.
Essentially the method of using nanoparticles brings improvements in properties of
epoxy resin and this development has taken centre stage of research in the recent
years.
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In order to widen the application areas of epoxy resin, it is important to
improve many of the epoxy resins properties. Therefore, researchers have started
looking at new methods of epoxy resin modification. Because of its special structures
and properties, the nanoparticles have become one of the most possible choice for
such modifications. In recent years, a large quantum of studies have been carried out
on the application of nanofiller filled epoxy composites and the research has shown
that epoxy resins with nanofillers have superior physical, mechanical, thermal and
dielectric properties.
1.2 Structure and properties of epoxy resin
Epoxy resin belongs to the epoxy oligomer class. It is able to react with a curing agent
or hardener to form a three-dimensional network. The most important advantage of
epoxy resin is that by choosing different epoxy oligomer and hardener or by suitable
modification methods, many improvements in material properties can be achieved.
Therefore, it has become one of the most popular thermoset polymer for industrial
applications.
1.2.1 Definition and classification of epoxy resin
Rings formed by two carbon atoms and one oxygen atom are called epoxy or epoxy
group. A compound that contains such rings are called as epoxides. The simplest
epoxide compound is ethylene epoxide, which is able to form thermosetting
polyethylene oxide through ionic polymerization. Such polyethylene oxide is known
as epoxy resin.
Epoxy resin is the collective name for compounds that contain two or more
epoxy groups mixed molecules and is able to form three-dimensional net structure
solidifying under chemical reagent [2]. In order to distinguish them from solidified
products, sometimes epoxy resins are also called epoxy oligomer because of
molecular weight belongs to oligomer. The main characteristics of epoxy resin
chemical structure is that there are epoxy groups with epoxy molecular chain.
However, by using different raw materials and methods of synthesis different
characteristics are achieved.
Epoxy resins can be divided into glycidyl epoxy resins and non-glycidyl
epoxy resins, according to the method of synthesis. The glycidyl epoxy resins are
formed by a condensation reaction of appropriate dihydroxy compound, dibasic acid
or a diamine and epichlorohydrin, whereas the non-glycidyl epoxy resin are prepared
by peroxidation of olefinic double bond. According to the synthesis methods, glycidyl
ether resin, glycidyl ester resin and glycidyl amine resin belong to glycidyl resins,
whereas alicyclic epoxy resin and aliphatic epoxy resin are classified into non-
glycidyl epoxy resin.
Further, glycidyl ether epoxy resin can be divided into diglycidyl ether of
bisphenol-A (DGEBA) and novolac epoxy resin. Both of DGEBA and novolac epoxy
resins are the most commonly used resins in industrial applications. DGEBA epoxy
resins are produced from reactions between epichlorohydrin and bisphenol-A. The
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characteristic property of molecular structure of this epoxy resin is that the molecular
chain contains active epoxy groups. Because there are active epoxy groups in the
molecular chain, the epoxy resin is able to support cross-linking reactions with
hardeners to form 3D cross-linked polymer net.
Novolac epoxy resins are formed under a reaction between phenolic novolac
resin and epichlorohydrin. In Comparison to DGEBA, novolac epoxy resins contain
more than two epoxy groups in their molecular structure. Therefore the cured
products have larger cross-linking density, better thermal stability, mechanical
properties, dielectric properties, water and corrosion resistance.
Glycidyl ester resins have better dielectric properties and weather resistance.
Their viscosity is normally lower than other epoxy resins. Glycidyl ester resins have
better adhesion characteristics than other epoxy resins. Moreover, they maintain their
properties at very low temperatures, which mean that their adhesive strength is still
higher than other epoxy resins at very low temperature. Glycidyl amine resins, on the
other hand, have high epoxide equivalent, big cross-linking density and higher
thermal resistance and hence they are preferred in carbon fiber reinforced systems.
Alicyclic epoxy resins are prepared by epoxidation of alicyclic alkene’s
double bond and their molecular structure has a large difference in comparison to
DGEBA and other epoxy resins [1]. This is because the epoxy groups of DGEBA and
other epoxy resins are directly connected by their alicyclic ring, whereas alicyclic
epoxy resin epoxy groups are connected by aliphatic hydrocarbon or benzene
molecules. Alicyclic epoxy resin curing products have high compressive and tensile
strength. Moreover, the alicyclic epoxy resin can maintain good mechanical properties
even at higher temperatures for long service periods.
1.2.2 Curing of epoxy resin
The epoxy resins before curing are only sticky liquids but do not have any practical
value. It needs to be cured into a three-dimensional cross-linked network structure
before being put into use. Such curing processes involve reaction between the epoxy
group and the curing agent (also called hardener). This reaction between epoxy
groups and hardeners is able to form a three-dimensional cross-linking network
structure and therefore the epoxy resin is able to cure into solid materials which are
firm and infusible.
Epoxy resins are always easily cured under certain conditions with curing
agents. Most of the epoxy resins, for example, bisphenol-A type epoxy resin, have
strong temperature stabilities. The bisphenol-A type epoxy resin is able to remain
structurally unchanged up to 200°C. However, epoxy resins also have a strong
reactivity as well. Thus they are able to react in presence of certain curing agents.
However, different epoxy resins have different curing requirements. Some epoxy
resins are able to cure under low temperature or room temperature, whereas others
require curing under high temperatures.
As the curing agent is able to influence epoxy resins properties significantly, it
also necessary to consider the effects of curing agents used. There are many types of
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curing agents and therefore it is necessary to choose curing agents according to the
requirements of application.
The curing temperature is also one of the important parameters that need to be
considered when selecting hardeners. Because the curing process is basically a
chemical reaction, any increase in curing temperature will increase the reaction rate
and will reduce the time taken for curing. However, it is also necessary to note that if
the curing temperature is too high, the epoxy resin may be cured unevenly. The
resultant cross-linking density will be asymmetric and may affect the ultimate
properties. Therefore, it is necessary to consider the curing temperature’s upper limit
during the curing process and maintain a balance between the curing time and desired
properties of the end product.
1.3 Nanocomposite materials
The nanomaterials have always shown unique optical, mechanical, thermal, magnetic
and electrical properties which are different from ordinary materials. In 1984 German
Scientist H. Gleiter successfully produced nanosize metal powders. Following this,
the nanometer sized materials were also introduced [3]. Because of their extremely
different properties as compared to the ordinary materials, nanometer size materials
and nanostructures have become the most attractive area of R&D in advanced
materials.
The nanoparticle reinforced epoxy resins have shown huge improvements [4,
5] and scientists believe that the improvements of epoxy resin properties are the result
of nanosize particle surface effects, quantum size effect and macroscopic quantum
tunneling effects (MQT) [6,7]. Because of the high viscidity of epoxy resins, it is hard
to mix nanosize fillers uniformly into epoxy resins. Hence it also necessary to
consider these aspects in manufacturing processes.
The word “nanocomposites” was suggested by Roy and Komarneni in 1984
[8]. A nanocomposite [9] is defined as “a multiphase solid material where one of the
phases has one, two or three dimensions of less than hundred nanometers (nm), or
structures having nanoscale repeat distances between the different phases that make
up the material”.
Since the inorganic nanoparticles have large surface area; the interface area
between the inorganic nanoparticles and polymers is large and the interface stresses
will reduce significantly. Hence the problem of unmatched thermal expansion co-
efficient of inorganic nanoparticles and the polymer base material has been solved.
Thus it is easy to make full use of the excellent mechanical properties and heat
resistance property of the in organic nanoparticles and the flexibility and processing
abilities of the polymer. The physical properties of polymer/inorganic nanocomposites
are much better than the ordinary composites. In recent years, many studies on
nanocomposites have yield significant results [10, 11].
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1.3.1 Particle dispersion and surface treatment
1.3.1.1 Particle dispersion
When the nanosize fillers are dispersed into the epoxy resin to form composites, the
nanoparticles tend to agglomerate with each other as a result of the interactions
between nanoparticles and dispersion is frequently observed in epoxy
nanocomposites. With the result, epoxy nanocomposites tend to lose their unique
characteristics as a result of agglomeration of nanoparticles. Thus it is a great
challenge to achieve a uniform and stable dispersion of nanoparticles in the base resin.
The Van der waals forces and the columbic forces between nanoparticles are partially
responsible for agglomeration. Such agglomerations caused by interaction of forces
can be minimized by application of mechanical forces and chemical reactions. The
presence of chemical bonds could also lead to particle agglomerations. To help the
development of polymer nanocomposites with consistent and improved quality, it is
necessary to understand the factors that influence the particle dispersion in the base
resin.
1.3.1.2 Surface treatment with silane solution
To obtain better particle dispersion, different methods based on mechanical and
chemical techniques have been developed to prevent agglomeration of nanoparticles.
The mechanical methods are intended to break the interaction force between
nanoparticles agglomerations and modify the surface structure of nanoparticles. The
most frequently used simple mechanical methods are high speed mixture technique
and ultrasonic dispersion method. The ultrasonic dispersion methods have become
more popular in recent years. By applying ultrasonic waves to the mixture, the
agglomeration of nanoparticles could be broken and a uniform dispersion of
nanoparticles can be obtained. However, it is also necessary to note that the surface
activity of nanoparticles could be increased by the use of high energy ultrasonic
waves. Long time exposure to ultrasonic waves leads to high possibilities of collision
between nanoparticles and this result in the formation of newer agglomerations.
The surface treatment on nanoparticles is another way to achieve better
particle dispersion in polymer nanocomposites. The surface treatment with silane as a
coupling agent is the most popular chemical method to modify the surface structure of
nanoSiO2, Al2O3 and ZnO particles it widely used. The silane coupling agent contains
functional groups that could react with both inorganic fillers and organic polymer
matrix. The silane coupling agent normally contains two types of groups, as shown in
Figure 1.1.
Figure 1.1. Chemical structure of silane coupling agent
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Silanol molecular chains are able to form oligomer structures and from
hydrogen bonds with the surface of inorganic nanosize fillers. Moreover, additional
condensation reactions will also occur between the coupling agent, Silanol groups and
the surface hydroxyls of inorganic fillers. Further, condensation and dehydration
reactions can also be obtained by drying nanofillers after surface treatment with
silane. The inorganic nanosize fillers with organic functional groups that are attached
to their surface by strong chemical bonds can be finally obtained as shown in Figure
1.2 [12].
Figure 1.2. Mechanism of silane surface treatment
The compatibility between nanoparticles with silane surface treatment and
polymer matrix is better compared with the nanoparticles without silane surface
treatment. The presence of hydrogen bonds increase the surface tension of inorganic
nanoparticles [13] and a more uniform dispersion of nanoparticles in the epoxy
nanocomposites can be achieved, as illustrated in Figure 1.3.
Figure 1.3. The dispersion of nanoparticles before and after silane surface treatment
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Recent studies have shown that the nanoparticle dispersion has strong effects
on the insulating behavior of resulting polymer nanocomposites. The dielectric
properties of polymer composites may be enhanced by surface treatment methods [14]
since the use of coupling agent will result in an increase in interfacial interactions.
There are many factors which may affect the effectiveness of surface treatment such
as the type of coupling agent chosen, the coupling agent concentrations, the duration
of treatment and the dispersion method. In order to obtain better dispersion, there is a
need to consider the process of surface treatment [15].
1.4 Interfacial characteristics
1.4.1 Inter-particle distance and surface area
The enhancement of insulating properties in polymer nanocomposites is mainly due to
the increases in specific surface area and the decrease in inter particle distance. In
polymer nanocomposites, by assuming that the spherical nanoparticles are uniformly
dispersed in the base material, the inter particle distance between any two
nanoparticles is proportional to particle diameter whereas the surface area is inversely
proportional to the particle diameter. Figure 1.4 shows the schematic diagram for
inter-particle distance and surface area calculation Both of the inter-particle distance
and the surface area are calculated by using the equation given by Tanaka and co-
workers [16].
Figure 1.4. Schematic diagram for inter-particle distance and surface area calculation
The inter-particle distance D as a function particle diameter d is described as, 1
321
6
f
D dV
(1.1)
where, Vf is the volume fraction of the nanoparticle in the base material. The
surface area per unit volume S of the particles is shown as,
6
fVS
d (1.2)
The above two formulae can also be expressed based on the weight percentage
of the nanoparticles, as shown below, 1
3100
1 1 16 100
n m
m n
wt%D d
wt%
(1.3)
d d D
8
22
3
1
d dS
D d D dD d
(1.4)
where, ρm and ρn are the specific gravity for the polymer matrix and the
nanoparticles respectively. Based on the equations 1.3 and 1.4, the inter particle
distance and the surface area of nanoparticles can be determined. For a typical epoxy
nanocomposite that contains 5wt.% nanoparticle with a diameter of 40nm, the inter
particle distance is 46.7nm and the surface area per unit volume is 10.8km2/m
3. In
comparison, the inter particle distance between 100µm particles is 117µm and the
surface area per unit volume is 0.0043km2/m
3. It can be seen that the nanoparticles
have extremely large surface area. It is necessary to consider the influence of
interfacial region between nanoparticles and polymer matrix on the dielectric
properties of nanocomposites as the area of interfacial region is associated with the
surface area of nanoparticle.
When nanoparticles are dispersed into polymer materials, the nanoparticles
tend to be in equilibrium with each other in the polymer nanocomposite due to the
forces of interaction. Thus interaction forces are essentially constant with each
particle although they might vary with the inter particle distance between the particles
[10,17]. The surrounding area of nanoparticles is increasingly modified by the
nanoparticles. Those surrounding areas that have different forces in comparison to the
base polymer materials are defined as the interface between nanoparticles and the
polymer matrix. The interface regions have mechanical, chemical, thermal and
insulating properties which are different from bulk materials and have significant
influence on the overall properties of the polymer composites [18]. Moreover,
particles with smaller size tend to have thicker interface region resulting in a
significant effect on the resulting nanocomposites. The specific surface area of
particles increases sharply as the diameters of these particles reduce to a sufficiently
low value. The high specific surface area results in a large interface region and high
surface defects.
In polymer nanocomposites, if the nanoparticles are uniformly dispersed in the
base polymer materials, the inter particle distance tends to be distributed according to
the Poisson distribution [19]. In this case, for spherical nanoparticles which are
uniformly distributed in polymer materials, the probability for the interface region
surrounding a nanoparticle to overlap with another nearby interface region is given by
the following equation,
21
tP exp
d (1.5)
where P is the probability of overlapping, t is the thickness of interface region
and d is the inter particle distance. For given filler loading concentration, the
probability of interface overlapping as a function of interface thickness over inter-
particle distance is shown in Figure 1.5. Thus for nanoparticles with an average inter-
particle distance of 40nm, there is a 50% possibility for its interface region to overlap
with a nearby interface if the thickness of interface is 14nm.
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Figure 1.5. Probability of interface region overlapping
1.4.2 Models for polymer nanocomposites
1.4.2.1 Dual layer model
In the interfacial region the properties of nanoparticles and the polymer matrix differ
due to the interaction between the two phases. Hence, a dual layer model was
proposed by Tsagarapoulos and co-authors [20] to help in understanding the interface
behavior of polymer nanocomposites. A schematic diagram of the dual layer model is
shown in Figure 1.6. In this model the interfacial area between nanoparticles and the
polymer matrix is divided into two different layers [20, 21]. The inner layer that
surrounds the surface of nanoparticles is assumed to be a tightly bond layer where the
polymer chains are tightly bonded to the surface of nanoparticles and those polymer
chains which are highly restricted. There is also another layer that surrounds the inner
layer where the polymer chains are loosely bound. The outer layer is named the
“loosely bound layer” and the thickness of this layer is slightly greater than the tightly
bound layer. The polymer chains tend to have higher mobility in the loosely bound
layer. It is also easier for charge carriers to move in the loosely bound layer.
Figure 1.6. Schematic diagram of dual layer model
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1.4.2.2 Multi-core model
Based on the idea discussed by Lewis [22], Tanaka proposed a multi-core model, as
illustrated in Figure 1.7. In this model it is assumed the spherical nanoparticles are
uniformly distributed in the base polymer materials, the interface area between
nanoparticles and polymer matrix can be classified into three different layers [16].
The first layer which is closest to the surface of nanoparticles is a bounded layer that
tightly bound to the surface of nanoparticles and polymer by coupling agents such as
silane. The second layer is a bound layer that contains polymer chains which are
strongly bound to the first layer and the nanoparticle surface. This bound layer is
more tightly bound as compared to the base polymer matrix.
The third layer is a loose layer which is loosely bound to the second layer.
Because of the surface tension effects, the loose layer has high free volume as
compared to base polymer matrix. Moreover, an electric double layer, which is known
as Gouy-Chapman diffuse layer, is also formed in the interface region. Due to the
effect of this electric double layer, the charge carriers with opposite signs are diffused
outwards from the interface region to the Debye shielding length.
The multi-core model provides an understanding of many dielectric
phenomena observed in polymer composites. Generally speaking, the bound layer and
the loose layer are the main factors that affect the dielectric performance of the
polymer composites. The presence of both bonded layer and the bound layer is
believed to restrict the mobility of polymer chains and leads to an increase in
dielectric constant and the glass transition temperature. On the other hand, the third
layer with large free volume is responsible for the reduction in both dielectric constant
and the glass transition temperature. The increase in mobility observed in some
studies is also believed to be due to the presence of shallower traps in the loose layer.
Figure 1.7. Schematic diagram of multi-core model
With the above discussion, it is clear that many properties and models are to be
considered for good understanding of the behavior of nanocomposites. The aim,
objectives, and scope of work are based on the discussions presented above. To
further supplement this background information, results of published literature are
discussed in the next section.
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1.5 Dielectric properties, polarization and depolarization of epoxy
nanocomposites
1.5.1 Principles of dielectric spectroscopy
Dielectric spectroscopy depends on the polarization that is induced in a material due
to the effect of an external electrical stress. Dielectric spectroscopy can provide useful
information on the electrical properties of the specimens. Moreover, this technique
can be used as analytical tool whereby the dielectric data is related to other properties
such as the polymer molecule structure or morphology [23], or its degradation and
ageing.
The interaction of electric field with matter is of fundamental importance in
basic and applied science. Many aspects of dielectric response, especially in the
presence of non-linear processes, are more easily understood in terms of the response
to time dependent signals [24]. However, there exists a very powerful alternative
approach which offers a very considerable theoretical and practical advantage,
provided that a linear system is studied and this is the determination of the response to
harmonic excitation, i.e., sinusoidal waves [24].
The mathematical basis for the treatment of the frequency domain response
rests on the Fourier transformation of a given function of time G(t), defined by the
Fourier transform [24],
1
22
F G( t ) ( ) ( ) G( t )exp( i t )dt (1.6)
The Fourier transform gives the frequency spectrum ψ (ω) of the time-
dependent function G (t) - the amplitudes, phases and frequencies of the sinusoidal
waves which make up the given time signal. The inclusion of i in the above
expression means that any transformed term will be complex, i.e., it will posses both
real and imaginary components. This indicates that any resulting expressions will take
into account the phase behavior of the response of a specimen [23] as well as the ratio
of the amplitudes.
1.5.1.1 Response of dielectrics to electric fields
The interaction of electromagnetic fields with matter is described by Maxwell’s
equations,
E B
t (1.7)
H j D
t (1.8)
0divB (1.9)
edivD (1.10)
E and H describe the electric and magnetic fields, D describes the dielectric
displacement, B, the magnetic induction, j the current density and ρe the density of
charges [25]. For small electric field strengths D can be expressed by, *
0D E (1.11)
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where, εo is the dielectric permittivity of vacuum (εo = 8.854×10-12
F/m). ε* is
the complex relative dielectric permittivity [26]. For a periodic electrical field E (t)
=Eo exp (−iωt) ω is the radial frequency, 𝑖 = −1 ) the complex dielectric function
ε* is defined by,
' "( ) ( ) ( ) i (1.12)
where, ε '(ω) is the real part and ε "(ω) the imaginary part of the complex
dielectric function. The polarization P describes the dielectric displacement which
originates from the response of a material to an external field only. Hence it is defined
as,
0 0 0 0( 1) P D D E E N E (1.13)
With χ * = (ε
*−1), where χ
* is the dielectric susceptibility of the material under
the influence of an external electric field. N0 is the number of dipoles per volume unit,
α is the polarisability of the charge. The equation also connects the dielectric
displacement with the contributions from the geometrical and polarisability of the
material.
1.5.2 Types of polarization
The polarization behavior is directly related to the electronic, atomic, orientational
and interfacial polarization of the material. The first two of these are induced by the
applied field and are caused by displacement of the electrons within the atom (its
polarization time scale is ~10-15
s) and atoms within the molecule (~10-12
to 10-13
s),
respectively. The orientational polarization (~10-9
to 10-3
s) is the classical type of
polarization originally treated by Debye and only exists in polar materials, i.e., those
with molecules having a permanent dipole moment [27]. There are therefore no
restoring forces tending to impose a preferred direction, except randomizing influence
of thermal agitation. The interfacial polarization appears in heterogeneous materials in
which the relaxation time is longer than that of the orientational polarization [28].
It comes from the accumulation of charges at the interfaces between the
various phases constituting materials when these various phases have different
dielectric constants and conductivities. For very conductive solutions, a layer of ions
will form adjacent to the electrodes [29-32]. This will alter the charge distribution
within the system and results in a marked rise in capacitance as the frequency is
lowered [23]. This effect is known as “electrode polarization” that normally is an
unwanted effect and should be removed or corrected for. It is possible that some of
the effects observed around nanoparticles in the composites studied here result in
similar effects.
Electronic and atomic polarization are temperature independent, but
orientational polarization, depending on the extent to which the applied field can
order the permanent dipoles against the disordering effect of the thermal energy of
their environment, varies inversely with absolute temperature. All of these
polarization mechanisms can only operate up to a limiting frequency, after which a
further frequency increase will result in their disappearance. Because of the spring-
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like nature of the forces involved, it is possible to observe the specific polarization at
a specific frequency region [24].
1.6 Dielectric response in the frequency domain
Dielectric analysis usually involves applying a field of fixed or varying frequency to a
specimen and measuring the response. As the frequency of the field changes, different
mechanisms of polarization will predominate. It is the analysis of these mechanisms
that provides the basis of dielectric spectroscopy [23]. Since some polarizations are
temperature dependent, the activation energies and dipolar types of some specific
relaxation processes can be obtained by measuring the dielectric responses over a
range of temperatures.
When an external field is removed from a charged capacitor, that capacitor
will discharge energy stored over a period of time, depending on its capacitance and
the resistance in the discharge circuit. In an alternating system, charge movement will
change direction in order to “keep up” with the fluctuations in the field when that
field charges direction. As this realignment will inevitably be non-instantaneous, the
response will take place over a period of time. According to equation 1.6, the
frequency dependent response function can be described in the time dependent
response function by Fourier transform.
If the Fourier transform is applied to polarization phenomena, then * * *
0 ( ) ( )P i E (1.14)
where χ is the susceptibility of the specimen, which is complex and may
therefore be expressed in terms of its real and imaginary components, i.e., *( ) '( ) "( ) (1.15)
The susceptibility is related to the dielectric constant, which may also be
expressed in terms of the real and imaginary components, i.e. equation 1.12. The
results of capacitance C* can also be used to analyze the dielectric properties of the
specimen even if the dimensions of the specimen are unknown. As to the parallel
plate specimens,
* 4 4 4 4( ) '( ) "( ) '( ) "( ) '( ) 1 "( ) 1
A A A AC C iC i i
d d d d
(1.16)
where, A is the area of the specimen and d is the thickness of the specimen.
Maity and co-workers [33] have studied the dielectric spectroscopy results in
the frequency range of 10-3
Hz to 103Hz and temperature range of 25°C to 90°C to
characterize pure epoxy and epoxy nanocomposites prepared with as-received and
pre-processed Al2O3 nanoparticles. They observed LFD (Low Frequency Dispersion)
below 100Hz and the inclusion of nanoparticles lowers the effective real and
imaginary permittivity of the composite material. The cross-over frequency for pure
epoxy occurs at 8.66Hz and 49.75Hz at 25°C and 90°C respectively. The evaluated
activation energy of base epoxy resin is 0.26eV, however, it is 0.41eV for the 1% as-
received nanofilled composite. They concluded that the dielectric behavior of epoxy
resin is significantly affected by the incorporation of alumina nanofillers, at low
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temperatures. At higher temperatures, it is essential to functionalize the particles
before use.
Nelson and co-authors [34] have reported dielectric spectroscopy
measurement from 10-3
Hz to 106Hz for cured epoxy, uncured epoxy, nano and micro
TiO2 loading of 10wt.% in epoxy matrix with temperature range from 25°C to 115°C.
They concluded that, at 115°C and in low frequency region the real and imaginary
permittivities of the nanomaterial are parallel to each other and attributed this to LFD
behavior. At all the loading levels of 40wt.% the nanocomposite have slightly lower
bulk conductivity than the microcomposite. FTIR results are discussed based on the
bonding between the nanoparticles and epoxy chains and correlated to dielectric
properties.
Castellon and co-authors [35] have studied the dielectric behavior of epoxy
based compounds containing micrometric and nanometric silica using dielectric
spectroscopy in the range of 10-2
Hz to 106Hz and they have used thermal step method
for space charge accumulation. They have highlighted the fact that polar nature of the
dielectric affects the space charge accumulation in the samples with different
concentrations of silica filler. Higher values of dielectric constant obtained for
micromaterial as compared to nanometric samples at low frequencies is due to the
polar nature of material and micro fillers are more conductive as compared to pure
epoxy and results obtained for dielectric loss factor is also higher.
Fothergill and co-workers [36] have studied cross-linked polyethylene
(XLPE), epoxy and epoxy-glass nanocomposite systems for “sub-hertz” responses
from 10-3
Hz to 106Hz over temperature range of 20°C to 100°C. The authors have
reported that XLPE cable samples at low temperatures exhibit percolation like
deterministic-fractal-circuit model described by Dissado and Hill [37] and epoxy
samples under thermal ageing exhibit QDC behavior. They also considered the effects
of water on the epoxy-glass nanocomposites. For water absorption, the real
capacitance in epoxy-glass composites does not increase much at 0.001Hz except at
higher humidity levels for 48h and they have suggested that this is due to charge
percolation.
Plesa and co-workers [38] have measured the relative permittivity (r') and the
loss tangent of epoxy resin with and without inorganic nanofillers alumina (Al2O3),
silica (SiO2) and titania (TiO2) using dielectric spectroscopy. The measurements are
carried out in the frequency range of 10-3
Hz to 106Hz and at temperature at 27°C and
60°C. They concluded that, the frequency dependence of the dielectric properties
emphasized low frequency dispersion for both unfilled and filled epoxy samples and
explained it in terms of intracluster and intercluster charge motion. This has little
influence on the type of the filler at low filler concentration. They have observed that
the dominant effect on the dielectric behavior is not the filler permittivity, but the
filler-polymer interface. With low content of fillers (1wt.%) in epoxy, the dielectric
behavior neither improves nor worsened. However, at this filler loading mechanical
and/or thermal properties show improvements. The increase of the TiO2 filler
concentration up to 5wt.% leads to lower r' values with respect to the unfilled epoxy.
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Frechette and co-authors [39] have investigated and compared the dielectric
spectroscopy results on epoxy-silica nanocomposites and generic compound
containing 60wt.% micrometric silica and a few % of nanosilica. They concluded
that, post treatment using heat was necessary to obtain reproducible dielectric
response. With a high content of micrometric silica, the dielectric response
(permittivity and losses) is dominated by the micrometric phase, but when the silica
content is kept invariant and microphase is replaced by a nanophase, the dielectric
response is different. When the nanophase reaches 5wt.%, its dielectric response
resembled that of the microcomposite. A large content of nanofiller in presence of
highly loaded microcomposite results in poor interfaces. In the present context, the
optimum nanofiller loading would appear to be between 2.5 and 5wt.%. They also
observed, dielectric features linked to the material compositions were annotated. For
the case of 65% total weight of silica, it was found that when the nanophase varied
from 0 to 5wt.%, the increase in the dielectric constant could fluctuate by as much as
10% in the low-frequency range.
The effect of glass-transition behaviors of silica, silver, aluminum, and carbon
black epoxy nanocomposites with their counterparts of micrometer sized fillers have
been investigated by Sun and co-authors [40]. It is reported that epoxy
nanocomposites show a Tg depression and suggested the reason for decrease in Tg
from thermo mechanical and dielectric relaxation processes, adsorbed water and
bonded organics at the surface of the nanosilica assisted the polymer relaxation
process at the filler-resin interface. However, the adsorbed water increases the
dielectric loss at low-frequency range. It is further observed that from dynamic
mechanical properties characterization, nanosized silica reduced Tg of the composites
but did not influence the sub-Tg transition temperature. Finally, they concluded that
the surface chemistry of the nanofillers and the interaction at the filler-resin interface
determines Tg and dielectric behavior of the nanocomposites.
Sun and co-workers [41] have studied the thermal properties, moisture
absorption and dielectric properties of pure epoxy and epoxy composites with micron-
sized silica and nanosized silica particles. They reported that the nanocomposite had a
much higher loss factor, lower glass transition temperature and higher moisture
absorption than micron composite and pure epoxy. They have presented theory and
models which focuses on the ionic contribution. They suggested that water influence
on the loss factor and the relaxation temperature of the nanocomposite is lower due to
the extra free volume at the filler-resin interface. They concluded that the dielectric
loss depends on the effect of moisture in epoxy and epoxy composite.
Yuang and co-workers [42] have reported the influence of the surface
treatment of silica nanoparticles, morphology, frequency and temperature dependence
of electrical conductivity, dielectric loss tangent, dielectric strength and dielectric
constant of both pure epoxy and the epoxy composites. Treatment of silane onto the
surface of silica nanoparticles improves the dispersion of the nanoparticles in epoxy
and electrical properties as well in comparison to untreated nanoparticles. Addition of
treated silica to epoxy resin increases the volume resistivity and decreases the
dielectric loss tangent over the temperature range above 77°C, but addition of
16
untreated silica decreases the volume resistivity and increases the dielectric loss
tangent over the whole temperature range of 27°C to 217°C. From the Weibull
distribution the characteristic, values of breakdown strength were calculated to be
30.8, 29.7 and 33.0kV/mm respectively for the pure epoxy, untreated and treated
silica filled epoxy.
Yang and co-authors [43] have studied and compared the effects of particle
size, dispersion on dielectric properties of epoxy-ZnO nanocomposites (NEP), epoxy-
ZnO micro composites (MEP) and deliberately not well dispersed nanoZnO
(NDNEP). They reported that at a loading of 5wt%, the three epoxy composites seem
to have no significant difference on resistivity as compared to epoxy resin; Dielectric
constants of all the epoxy composites are also basically the same but they are higher
than the pure epoxy samples. Dissipation factor (tanδ) of not well dispersed nanoZnO
is greater than that of epoxy-ZnO nanocomposites and epoxy-ZnO micro composites
and epoxy-ZnO nanocomposites has the minimum dielectric loss factor, whereas
dielectric loss factors of the three epoxy composites are larger than that of the pure
epoxy resin. The decreasing order of electrical breakdown strength for the three epoxy
composites and for the pure epoxy resin is as follows: NEP>MEP>NDNEP>EP. They
have proposed an aggregation interface phase model to explain the experimental
results.
Rouyre and co-workers [44] have studied electrical and mechanical properties
of epoxy polymers reinforced with a premixed epoxy/nanosilica (Nanopox F400)
master batch, dry non-surface-treated nanosilica powder and micro-silica flour. Cured
plates of epoxy containing both nanoparticles (up to 15wt.%) and microparticles (up
to 51wt.%) of silica were fabricated and evaluated. They have observed that the
tensile modulus increases with the addition of silica and experimental results agree
with data from theoretical models. Influences on dielectric permittivity, resistivity and
dielectric loss is observed to depend on the type of filler. High values of dielectric
strength were observed with micro-particles, while nanosilica had a slight negative
effect. Nanopox F400 silica fillers show a better dispersion in comparison with dry
powder, which has a direct consequence on dielectric permittivity (4.21) and
dielectric loss (0.0040) at 15wt.%. The measured glass transition temperature was
always close to 140°C. The fracture energy and electrical resistivity showed
improvements with microsilica and dry nanosilica.
Fleming and co-workers [45] have examined the conductivity over
temperature range of 30°C to 70°C and ac impedance from frequency range of 103Hz-
106Hz in air and in vacuum on samples of low density polyethylene to which
nanosized and microsized ZnO particles and a dispersant are added. The space charge
profiles were obtained using the laser-intensity-modulation-method (LIMM). They
conclude that, the temperature dependence of the vacuum dc conductivity in samples
containing the dispersant and 10% w/w nanosized ZnO, the conductivity shows
decrease of 1-2 orders of magnitude than that of a sample containing dispersant only
with temperature. But, addition of nanoparticles increased the ac conductivity at
higher frequencies. It is also observed that relative permittivity of samples with
nanoparticles also increases relative to that of samples containing dispersant only due
17
to homogeneous dispersion, at all temperatures, but the corresponding values in
samples with micro particles are unchanged. Space charge densities of order 300cm-3
are measured in the bulk near the electrodes and addition of nanoparticles slightly
decreases the density of homocharge which accumulate close to the electrodes.
1.7 Polarization and depolarization current
Polymer composites have excellent mechanical, dielectric and charge storage
properties. Therefore, in addition to mechanical and dielectric properties, it is
extremely important to understand charge storage characteristics of polymer
composites. Some of these characteristics are determined through frequency domain
spectroscopy (FDS) and polarization and depolarization current (PDC) measurements
[46]. Polarization current provides information about conduction and polarization
mechanisms [47] and therefore the effect of nanosized particle additives on properties
of composites has evoked considerable interest in the last decade [48].
Patel and co-authors [49] have studied the effect of nanofillers such as
spherical alumina (Al2O3), titania (TiO2) and zinc oxide (ZnO) particles and pre-
processing techniques on polarization and depolarization currents by varying
electrification time, temperature, electrode material on epoxy-based composites. They
have reported that, ZnO filled composites have low absorption current and low dc
conductivity followed by TiO2 and Al2O3 filled composites. They observe that
functionalization of particles before preparation of composites decreases the
conduction current in comparison to unfunctionalized fillers. The incorporation of
nanoparticles into the epoxy increases the activation energy in alumina
nanocomposites and at higher temperatures, polarization currents take less time to
stabilize. Current variations are similar in trend at lower temperatures for both
aluminum and brass electrodes but at higher temperatures, but with brass electrode,
the current takes longer time to stabilize.
Li, Zhe and co-authors [50] have developed nano/micro-SiO2 dispersed low-
density polyethylene (LDPE) using double solution mixture method. They have
reported that the depolarization intensity of composite containing nano/micro
inorganic filler is lower than that of pure LDPE. The value of dielectric loss of pure
PE is higher than that of the composite during depolarization processes. At higher
temperature (333K), the peak of dielectric loss is highest at 1wt.% of nanoSiO2 as
compared to 3 and 5wt.% filler loading of SiO2. They have concluded that
polarization behavior has a close relation to the physical characteristic of fillers, as
well as the filler content.
Yin and co-workers [51] have investigated depolarization currents using
thermally-stimulated current (TSC) for the composite samples of low-density
polyethylene (LDPE)/ nanoSiO2 and LDPE/microSiO2 which are prepared using
double-solution mixture method. An initial rise method is used to investigate
activation energy and depolarization current of both composites and pure LDPE. They
have shown that, LDPE has the greatest activation energy among all the samples,
whereas, the activation energy of the interfaces between nanoparticles is shallower
18
than that of amorphous and spherocrystal structures and trap depth decreases with
nanoSiO2 loading from 0.5 - 3.0wt.%. This phenomenon is discussed on the basis of
multi-core model [16]. Further, the activation energy of the interfaces between micro-
particles is slightly shallower than that of amorphous and spherocrystal structures and
does not vary with the loading of micro-particles.
Xingyi and others [52] have studied the effects of the surface modification of
Al nanoparticle fillers on the electrical conduction behavior of PE nanocomposites by
means of polarization and depolarization current of PE and PE loaded with un-
surface-treated and OTMS (Octyltrimethoxysilane)-coated Al nanoparticles. They
observed that the incorporation of octyl groups onto the surface of Al nanoparticle
alters the time dependence behavior of polarization and depolarization current and
increases the percolation threshold and resistivity of the polyethylene composites.
They have concluded that based on morphology of the nanocomposites, improvement
in electrical properties may be achieved by good dispersion and surface-treatment of
nanoparticles in the polymer matrix.
Castellon and others [53] have analyzed electrical properties like charge
injection, polarization, trapping and conduction phenomena under multiple stresses
for epoxy based matrix which are filled with micro and/or nanoparticles of silica. The
Schottky Injection and Space Charge Limited Current models are studied to explain
the conduction phenomena. They have concluded that electrical properties of the
polymers are strongly influenced by micro silica content. Both polarization and
conduction phenomena show increase with silica content. Based on this result, they
have explained that there is weak space charge accumulation of the polymers at
(micro and nano) 65wt.% of silica content of micro and nanosizes. They have
proposed a composition of sample containing 62.5wt.% of micro and 2.5% of
nanofiller for achieving lower concentrations of space charges in the polymer.
Smith and co-authors [54] have carried experiments on absorption current on
12.5wt.% micro silica/XLPE, silica/XLPE and vinyl silane/XLPE nanocomposites at
an applied electric stress of 30kV/mm. They have observed that the nanocomposites
display a classic nI(t) At characteristics till charge front arrives at the electrode and
it takes place at ~500s for XLPE, while for the composite it occurs at 1000s, during
which there is a demonstrable change in the slope, indicating that charge mobility is
reduced by a factor of 2. After poling, calculation of the current decay exponent for
XLPE takes place at 1.34 decades of current for each time decade; this value is higher
for the microcomposite but is considerably lower for the nanocomposites. This is in
agreement with the scattering/reduced mobility hypothesis.
Chen Zou and co-workers [55] have studied the time dependent direct current
conduction characteristics of both micro and nanosilica filled epoxy composites at 0%
and 75% relative humidity in the temperature range of 25°C to 70°C. They have
concluded that, ohmic conduction dominates within dry epoxy at low electric field
until the buildup of space charge takes place. Water enhances the charge decay and
increases the threshold field of space charge accumulation in the epoxy materials. In a
humid environment, when more inorganic filler is added, more water is dispersed into
19
epoxy and it is more difficult to accumulate space charge within the epoxy. At low
filler loadings, composites have similar conduction mechanisms as the polymer
matrix, in which ohmic conduction and trap sites limit the conduction. But at higher
electric fields, the phenomena is dominant with increasing loading of nanoparticles,
since the barrier heights of the epoxy nanocomposites show tendency to decrease.
1.8 Electrical conductivity/resistivity
There is enough literature available on the aspects of electrical conductivity of
polymer composites with both thermoplastic and thermoset polymers and different
types of conducting fillers and experimental results show the existence of a
percolation threshold for every system [56-58]. A majority of the reported
experimental work were performed with carbonaceous fillers (mainly carbon black) in
different polymer systems. The percolation threshold depends upon the conductivity
of the fillers and the distribution of fillers in the composite [59]. The filler
conductivity in turn is influenced by the filler type and the filler structure, size and
shape [58, 59]. Apart from the filler properties, experiments further demonstrate that
polymer types also influence the percolation threshold and this effect can be observed
from the variations in the conductivity of different thermoplastic polymer composites
having the same carbon black filler [58]. Further, composite conductivity is also
influenced by the surface properties of the polymer and the filler due to the interaction
dynamics between the polymer and the filler [59].
Some amount of work has also been reported on the dc and ac non-linear
conductivity of zinc oxide (ZnO) filled polymer composites [60, 61]. The primary
idea was to develop an insulating polymer composite which would demonstrate surge
suppressive characteristics or which can be used for field grading applications.
Varlow and co-authors [61] have showed that with zinc oxide fillers in epoxy resin,
non-linear conductivity was observed under both ac and dc conditions when the filler
concentration exceeded the percolation threshold (~15%).
Tjong and Liang [62] have measured ac volume resistivity at 50Hz of LDPE-
ZnO nanocomposites and it is shown that a gradual decrease in the resistivity occurs
with increasing ZnO nanofiller loading of 60%, whereas, in the case of LDPE-ZnO
microcomposites, a percolation threshold at a filler loading of around 18% was
recorded and the variation in the trend in nanocomposites was attributed to the
differences in the inter-particle distance. In another study, Hong and co-workers [63]
have investigated the dc resistivity of LDPE-ZnO composite systems and reported
that the onset of percolation in LDPE-ZnO nanocomposites occurred at a lower filler
concentration as compared to the corresponding microcomposites.
Investigations by Cao and co-workers [64] have also showed that the volume
resistivity (ρv) of polyimide nanocomposites increases with the addition of nanofillers
at elevated temperatures and it has been reasoned that the behavior was due to the
charge trapping in the nanocomposites. Sarathi and co-workers [260] have reported
the effect of temperature on the composites was minimal and with water aging,
considerable reduction in ρv, irrespective of filler loading was observed. Further, for
20
epoxy-layered silicate nanocomposites system with 5wt.%, nanoclay, a decrease in ρv
with temperature as compared to pure epoxy was observed. It is suggested in the work
[65] that microsized TiO2 particles incorporated in LDPE increases electronic
injection from the electrode and also act as charge traps in the bulk.
Bulinski and co-workers [66] have reported that the nanocomposites when
subjected to a dc field at room temperature and aging at 90°C, up to 500h results in
the increase of conductivity. However, but the increments are significantly smaller
than those observed in the materials without organoclay. Based on these published
results, it appears that there is some influence of nanofiller (inorganic type) on the
resistivity of polymer nanocomposites.
1.9 Dielectric breakdown in nanodielectric composites
The dielectric strength of polymer nanocomposites has been generating tremendous
interest amongst researchers and several studies are reported in literature on this topic.
Ding and Varlow [67] performed electrical treeing breakdown distance
measurements under ac stress for epoxy-ZnO nanocomposites and observed that
addition of few weight fractions (<1%) of nanoZnO to epoxy led to significant
enhancement in the time to breakdown for the nanocomposite. Imai and co-workers
[68], have reported an increase of two times the electrical breakdown time for the
epoxy- layered silicate nanocomposites (with 5% filler concentration) as compared to
base epoxy at 20°C whereas at 80°C, the increase was observed to be around 6 times.
Nelson and Hu [69] carried out voltage endurance studies on epoxy-TiO2
composite systems and observed that nanocomposites have improved endurance
characteristics as compared to microcomposites, especially in the low electric field
regions. This improved voltage endurance for epoxy-TiO2 nanocomposites is
corroborated by Imai and co-authors [70] work which further demonstrated that epoxy
nanocomposites with SiO2 has enhanced insulation breakdown time as compared to
unfilled epoxy. Similar enhancement in the electrical breakdown time has also been
obtained for epoxy nanocomposites filled with Al2O3, nanoparticles at a filler loading
of 5% filler. In another study, Roy and co-workers [71] have examined the influence
of treated fillers on the voltage endurance properties of XLPE-SiO2 nanocomposites
and they observe that the nanocomposites containing treated fillers exhibit higher
electrical endurance.
In contrast to the observations made for the voltage endurance studies for
polymer nanocomposites, results for the ramp type of voltage application under both
ac and dc stresses have shown a mixed trend. In some cases, an increase in the
dielectric strength of the nanocomposites is reported in comparison to unfilled
polymers and microcomposites whereas in few other reports, a decrease is reported.
Experimental studies by Cao and co-authors [72] have showed that the incorporation
of nanoparticles of alumina and silica into polyimide increases breakdown voltage of
the nanocomposite by around 10-15% up to 10% filler concentration. Similar
enhancements in the dielectric strengths have also been observed by Tuncer and co-
workers [73] for poly vinyl alcohol (PVA)-TiO2 nanoparticles under AC voltage and
21
by Murata and co-workers [74] for LDPE-MgO nanocomposites under dc voltage.
The occurrence of higher electrical breakdown strengths of nanocomposites is further
supported by the results of Hu and co-workers [75] have shown that the dc, ac and
impulse dielectric strength of epoxy-TiO2 nanocomposites are higher than that of base
epoxy and microcomposites. Contrary to the observations by Hu [75], Fuse and co-
workers [265] have shown that the dielectric strength of polyamide-mica
nanocomposites under ac, dc and impulse voltages are independent of the nanofiller
content up to 5% by weight. In another study, Nelson and co-workers [271] have
observed that although the dc electric strength values of epoxy-TiO2 nanocomposites
are lower than that of unfilled epoxy and the values are significantly higher when
compared to those of the microcomposites at a filler loading of 10%.
In a somewhat similar observation, Hong and co-authors [76] have measured
the dc breakdown strength of LDPE/ZnO nanocomposites and observed that the
values decreased with increasing filler concentration and higher filler loadings only,
and the breakdown strength of nanocomposites are higher than that of
microcomposites of the same materials. Contrasting dc dielectric strength behaviors
of polymer nanocomposites are reported by Zilg and co-workers [77] and they
observe that when organically modified layered silicates are added to EVA, the
electrical strength reduces whereas with addition of same filler PP, there is an increase
in the electrical strength. Imai and co-authors [70] have performed experiments to
measure the dielectric strengths of epoxy nanocomposites with different nanofillers
under homogenous and divergent ac electric fields. Results demonstrate that under
homogenous electric field conditions, the breakdown strengths of the nanocomposites
was almost the same or less than that of unfilled epoxy whereas for divergent field
conditions, nanocomposites had a higher electric strength as compared to base epoxy.
The influence of coated nanofillers on the dielectric strengths of polymer
nanocomposites under a ramp type of voltage application has been studied too.
Investigations by Ma and co-authors [78] have shown that when as received
TiO2 nanoparticles were introduced into LDPE, the dielectric strength of the
nanocomposite is less than that of unfilled LDPE whereas the same TiO2 with
AEAPS coating, display higher dielectric strength as compared to the unfilled
polymer. In another study on similar lines, Roy and co-authors [71] have performed
dc breakdown experiments at different temperatures on XLPE-silica nanocomposites
using nanosilica treated with three different compounds. Even when untreated silica
was used as the filler, the dielectric strength was observed to be higher than that of
unfilled XLPE and this increase is further enhanced when treated silica is used.
Surprisingly, the dielectric strengths of nanocomposites are still significantly higher
than that of unfilled polymer at higher temperatures of measurements. In another
interesting study, Imai and co-authors [79] have investigated the dielectric breakdown
characteristics of an epoxy composite containing a mixture of micrometer sized SiO2
and layered silicate nanofillers. The authors observed that the addition of just 1.5% of
layered silicate by volume to an epoxy-SiO2 microcomposite could increase the ac
dielectric strength of the nanocomposite to a value more than that of the base epoxy
composite. The voltage endurance of the same sample was also recorded to be higher
22
than that of the base composite. This result is significant since the properties of
traditionally used polymer microcomposites can be enhanced by the addition of small
quantities of nanomaterials.
The breakdown voltage is considered to be randomly variable, which
necessitates a statistical analysis of the electrical breakdown data. The Weibull
distribution [80] is on extreme value distribution and is most appropriate for such data
is well developed for the analysis of small and large data sets with censored data.
Other distributions used for electrical breakdown are the Gumbel and Lognormal [81].
Alternatively, Tuncer and co-workers [82] have proposed a different expression for
breakdown analysis.
The Weibull statistical distribution can be expressed as,
exp
o
EP
E
1 (1.17)
where, E is dielectric strength (V/μm), Eo (V/μm), known as the scale
parameter, is the electric field at which at least 63.2% of the samples are bound to fail.
The parameter β (dimensionless), known as the shape parameter, is a measure of
scatter in the data. A high value of β corresponds to lower scatter. For polymers, β
values in the range of 2-4 are commonly observed, as reported by Roy and co-authors
[258]. However, very high values for β are also reported in literature [83].
According to the recommendation of the IEEE 930-2004, a good, simple,
approximation for the most likely probability of failure is represented by the Equation
(1.18),
i
i .P %
n .
044100
025 (1.18)
where, i is the ith
result when the values of the dielectric strength (E) are sorted
in ascending order and n is the number of specimens. For this study, n = 20. The 90%
confidence intervals are calculated according to IEEE 930-2004 and the
corresponding “confidence interval tables” are obtained.
XingyiHuang and co-workers [42] have used the Weibull distribution to
analyze the breakdown strength data. For the pure epoxy, untreated silica and treated
silica filled epoxy the characteristic values of breakdown strength were calculated
statistically and found that the difference in breakdown strength between the epoxy
resin, and untreated silica filled epoxy are observed to be marginal. However, the
difference between the epoxy resin and treated silica filled epoxy are reported to be
statistically significant.
Fuse and co-authors [265] have studied nanocomposite system with polyamide
using layered silicate nanofillers from 1 to 5wt.%. They have reported that the
conduction current decreases with the addition of nanofillers and the dielectric
strength was almost independent of the nanofiller content for impulse, dc and ac
voltages. These researchers did not see an improvement in the use of nanofillers.
Many authors [84, 85] have reported that dielectric strength, as well as ρv and ρs, are
reduced by water uptake.
The damaging effect of metallic fillers on the dielectric strength of polymers
23
can be gauged from the drastic reduction in the ac dielectric strength of polyethylene
composites filled with aluminum particles at very low particle loadings [86].
Morshuis and co-authors [87] have performed experiments on polyethylene (PE)
cable samples and have shown that conducting inclusions are not as harmful as
mineral particles, in particular glass. They suggested that about 50% of the cables
polluted with glass failed during a type test for medium-voltage cables (10kV/mm,
24h) and all polluted cables showed a high failure percentage (>50%) at a type test for
high-voltage cables (23kV/mm, 24h).
1.10 Arc and tracking resistance in nanocomposites
Polymeric insulation is preferred because of better dielectric properties, low surface
energy which supports a good hydrophobic surface, better pollution performance in
outdoor service conditions [88]. In addition, polymeric insulators are light in weight,
easy to handle, vandal resistant and cost effective. However, the polymers have
certain disadvantages. The presence of moisture in polymers can cause chemical
hydrolysis, loss of plasticizer and the filler material causing embrittlement of material
or the pollutants get dissolved in moisture on the surface resulting in enhancement of
leakage current magnitudes and failure of insulation. There is an ever increasing
interest in power industry worldwide, to understand the degradation processes and the
performance characteristics of polymer insulating materials under severe pollution.
Arcing in insulating materials, either organic or inorganic brings about partial
disintegration of the surface layers of a material and changes its characteristics. Arc
resistance is the length of the time an insulating material can resist the action of an arc
while restoring its properties within a short time after arc extinction [89]. Arc
resistance is of great importance because of the origin of the arc is a hot point that can
burn insulating materials and cause flaming or create a conductive path.
Tracking is a phenomenon which occurs on the surface because of the reduced
creepage discharge resulting from surface contamination. Once tracking occurs, the
surface electrical insulation property is lost completely and it never recovers. In order
to improve reliability and performance of insulation materials, tracking phenomenon
is being investigated worldwide [90]. Guastavino and co-workers [91] have performed
tracking resistance evaluation of cycloaliphatic epoxy polymers with nanoclay and
microsilica. They suggested that 1.0 to 2.5wt.% nanoclay inclusion showed better
resistance to tracking and erosion under wet conditions. RajaPrabu and co-workers
[92] have studied tracking phenomenon in blends of silicone rubber and ethylene
Propylene Diene Monomer (EPDM). The test results show that the increasing
proportion of silicone rubber (100%) enhances the tracking and arcing resistance by
38% and 70% due to highly flexible bonding in silicone rubber.
Harindu and co-workers [93] have prepared HDPE and PBT blends and it was
observed that for all the compositions of HDPE/PBT compatibilized blends using E-
Ion1.150 (Na-ionic-graft-copolymer) and SPC (commercial) ionomer, resulted in
improved tracking resistance. It is also reported that the good results are obtained for
arc resistance and comparative tracking index for HDPE/PBT (80/20) using 3% of E-
24
Ion1.150 blend. This is due to the fact that ionomer rapidly reacts to form a cross-
linked structure with the ester group of PBT and the ethylene group of HDPE. This is
also attributed to the decrease in the rate of carbonization and more time is required to
form leakage current between the electrodes.
Tanaka and co-workers [94] have evaluated the tracking resistance of three
completely different polymeric systems, namely high temperature vulcanized silicone
rubber (HTV SiR), ethyl vinyl acetate (EVA) and epoxy resin. As observed, both the
HTV silicone rubber and EVA specimens are reinforced with alumina trihydrate
(ATH) for improving their tracking and erosion resistance properties, while the epoxy
resin compounds contain silane treated silica as reinforcing agent. A substantial
reduction in the tracking and erosion resistance of the polymeric materials is observed
with dc stresses in comparison to ac.
Park and co-workers [95] have developed nanocomposites with PTFE as base
material using different fillers. The arc resistance of PTFE composites was observed
to increase with increasing filler content, but values were different with different
fillers and their content. The arc resistance of Boron nitride (BN) filled PTFE
composites is better than that of Al2O3 and TiO2 filled PTFE composites. The most
important aspect of fillers seems to be the light reflectance under ultraviolet radiation.
Ratzke and co-authors [96] have demonstrated that nanofillers and microfillers
in a HTV silicone elastomer affect the resistance to arcing and the resistance to
tracking and erosion. They have reported that best dispersion is obtained for
nanosilica. On the other hand, large agglomerates are formed by nanoAl2O3.
Lei and co-authors [97] have studied two kinds of nanomaterials to modify the
properties of room temperature vulcanized silicone rubber (RTV SiR) under corona
discharge; These materials are, nanosilica and nanolayered silicate (at 2 and 5wt.%
concentrations). After corona aging, nanofilled RTV perform much better than the
virgin RTV material. It is suggested that nanofilled RTV has a superior corona aging
performance as compared to the virgin RTV material.
Formulations of RTV SiR with nanosilica as compared to RTV SiR with
micro silica have been evaluated by El-Hag and co-workers [98]. They have
suggested that the erosion resistance increases in direct proportion to the amount of
filler used. As a result, data on inclined plane tracking [98] confirmed that the
nanofilled SiR composites with as low as 10wt.% of nanofillers display a significant
improvement in resistance to erosion as compared with microfilled SiR composites.
Meyer and co-workers [99] have suggested that higher tracking and erosion
resistance, lower roughness, and slightly lower hydrophobicity are observed with
RTV SiR filled with nanosilica as compared to that of RTV SiR filled with micro
silica. The concentrations used in this work were 5 and 10wt.% for nano and micro
silica, respectively and nanosilica had higher tracking and erosion resistance as
compared to micro silica.
Polydimethylsiloxane (PDMS) incorporated with alumina trihydrate (ATH,
Al2O3.3H2O) filler is a representative system used in HTV-SiRs [100,101]. The ATH
is very efficient in enhancing the tracking resistance of insulating polymers [102,
103]. The resistance of epoxy resin to surface degradation can be considerably
25
improved with the addition of small volume fraction of spherical Al2O3 nanoparticles
[104]. The improvement in the resistance to surface degradation is further profound
when the nanoparticles are preprocessed before adding to the base polymer [104].
Sarathi and co-workers [260] have reported that the tracking time was higher
with epoxy nanocomposites as compared to virgin epoxy. With water aging at 30°C,
the epoxy nanocomposite at 5wt.% nanoclay showed decrease in tracking time by
32%. However, with thermal aging at 100°C for 360h, tracking time of epoxy
nanocomposites (5wt.%) is observed to be 52% higher as compared to virgin epoxy.
1.11 Water absorption and contact angle measurement in
nanocomposites
Many researchers have studied the contact angle to assess the hydrophobicity of
material surface. Furthermore, they have studied the causes of loss of hydrophobicity,
due to generation of hydrophilic groups, accumulation of surface charges and surface
crazing and cracking. These factors lead to increased surface roughness and erosion.
Measurement of the contact angle on polymer surfaces is a very useful method for
monitoring the migratory behavior of various organic functional groups from the bulk
to the surface [105]. A surface which has the inherent ability to maintain a high level
of hydrophobicity in the presence of moisture would have a low leakage current,
much reduced dry band arcing and therefore show better overall electrical
performance with longer lifetime. The measurement of the contact angle (degree) of a
water droplet on the polymer surface provides an indication of the state of the
hydrophobicity of the surface [106]. Hydrophobic materials allow less water surface
contact and thus make > 90°, whereas materials which are easily wettable allow
water to touch a large surface area and hence make < 90°. Surface is said to be
hydrophobic, when > 90°, hydrophilic when < 35° and partially wettable when
35° < < 90° [107].
Ethylene propylene diene monomer (EPDM) rubber is used as a construction
material for insulator, because it provides hydrophobicity for a longer time. The long-
term maintenance of the hydrophobicity is attributed to its chemical stability and
hydrophobicity recovery phenomena resulting from diffusion of low molecular weight
silicone polymers [108]. The hydrophobicity of polymers leads to higher electrical
surface resistance, but it is reduced because of water absorption during aging at
ambient temperature and contamination build-up.
A weakness caused by water may have a detrimental effect on the mechanical
[109] and electrical behavior of nanocomposites and it might nullify the effect of
nanofillers. If nanofillers are incorporated, the specific area of the epoxy-particle
interface is greatly increased as the particle dimensions decreases and this is a
potential location for ingress of water [110]. Water ingress into a polymer matrix
leads to different effects: plasticization through interaction of the ester molecules with
polar groups in the matrix, creation of micro crazes through environmental stress
cracking, leaching of unreacted monomer and in certain cases degradation of the
resin. Relatively short times of exposure lead to more or less reversible plasticization,
26
resulting in lowering of glass transition temperature [111].
The extent to which glass transition temperature is reduced depends on the
amount of water absorbed and is described by Kelly-Beuche equation based on the
free volume theory [112] or by the fox equation [113].
P P g W p g
g
P P W p
α V T (p) + α (1 - V )T (w)T =
α V α (1 - V ) (1.19)
p w
g g g
W W1= +
T T (p) T (w) (1.20)
where, Tg(p) and Tg(w) are respectively, the glass transition temperatures of
polymer and water, Vp is the volume fraction of the component, Wi is the weight
fraction and αi is the expansion coefficient; the subscripts p and w refer to polymer
and water, respectively. Agreement with Fox equation [113] implies that the water is
uniformly dispersed in the epoxy resin, rather than existing as isolated droplets or
clusters. The value of Tg can also be affected by hydrolysis of the cross-linked
polymer leading to a decrease in the cross-linking density [114, 115] and the
development of two phase structure [114].
Apicella and co-authors [116, 117] have reported that the amount of water,
taken up by an epoxy resin, depends on its hydrothermal history. Exposure to water
and then drying, leads to an increase in the equilibrium water uptake after each cycle.
It is assumed that the water causes irreversible damage to epoxy resins in the form of
micro-cavities and part of the water is molecularly dispersed in the polymer but a part
is left as resides in the micro-cavities. The subsequent growth of the cavities may
occur due to degradation of the resin matrix or extraction of residual by-products of
the synthesis of the resin. Sodium chloride is a common impurity in epoxy resins and
it is liberated as a result of reaction of epichloro-hydrin with the corresponding
phenol. When the voids contain an electrolyte, osmotic pressure becomes the driving
force for its growth into micro-cracks and micro-crazes [118]. Maxwell and Pethrick
[119] have used dielectric relaxation methods and it is reported that water might exist
as either clusters in voids or in a molecularly dispersed state within the resin matrix.
Antoon and co-workers [115] have suggested that water dispersed in the
epoxy matrix is usually strongly bonded to hydroxyl groups and its absorption is
completely reversible. However, Jelinski and co-workers [120] have stated that the
movement of water in epoxy resins is impeded and there is no free water and have
included that there was no evidence for tightly bound water. Woo and Piggott [121]
have suggested that in certain epoxy systems water is not bound to polar groups in the
resin or to hydrogen-bonding sites. It is clear that there are a number of unknown
factors with regard to the nature of the interactions between water and epoxy resins.
Long-term exposure to water can lead to loss of material by degradation due to
hydrolysis, oxidation/dehydration reactions involving loss of hydroxyl groups [122].
Comyn [118] has suggested that when epoxy resins are left in water for an
extended period, a small residue remains after evaporation of water, the origin of
which is not clear. Chen and co-workers [123] have studied the effect of water
absorption on the dielectric properties of epoxy, epoxy-micro composite and epoxy
27
nanocomposites filled with silica. They observed that nanofillers with unmodified
surfaces enhance the water sorption in epoxy materials, but micro fillers make little
contribution. The epoxy-particle interface is a potential location for water [123] and
the composites with very high specific areas may be particularly vulnerable to the
effect of water.
Hao and co-workers [124] have examined the water absorption of Epon
828/Epicure epoxy-clay nanocomposites and have suggested that some reduction in
water absorption due to the addition of nanoclay. Massam and Pinnavaia [125] have
investigated the resistance of epoxy nanocomposites towards organic solvents and
water. The absorption of methanol, ethanol and propanol in nanocomposites is
observed to be faster in pure epoxy systems and properties of pure systems are much
more affected by the absorbed solvent than nanocomposites. However, in the case of
water, only the rate of absorption is reduced. However, the equilibrium of water
uptake is relatively unaffected. It was further observed that the barrier to solvent
uptake is more significant in exfoliated composites than conventional or intercalated
layered silicate composites.
Becker and co-workers [126] have studied the water uptake of nanocomposites
by the direct mixing (DM) method, based on epoxy resins of three different
functionalities with 1.30E organoclay. The equilibrium water uptake at 80°C was
reduced by 4.76% for bi-functional (DGEBA) epoxy, 9.74% for tri-functional epoxy,
and 4.76% for tetra-functional (TGDDM) epoxy at 10% clay loading as compared to
the pure epoxy system. However, the maximum increases of diffusivity for these three
materials were 14%, 61% and 78% respectively. The concentration of clay did not
correlate proportionally with the reduction in equilibrium of water uptake or the
increase of diffusivity. This is attributed to the type of epoxy systems [126].
Numerous diffusion models have been proposed for modeling water
absorption in polymers and polymer composites. The most common approach is to
apply Fick’s law [127] to simple single-free-phase diffusion, due to its simplicity and
mathematical tractability [127]. However, it has been demonstrated that diffusion of
water in some glassy polymers is anomalous (non-Fickian) [128]. Two main
approaches are proposed to model the anomalous diffusion. One is the Langmuir-type
model for diffusion (LMD), assuming that absorbed water molecules consist of
mobile and bound phases [129, 130] and the other is the diffusion with time-varying
diffusivity model (DTVD), where constant conductivity-efficient diffusion is replaced
by a decreasing function of time analogous with a relaxation models of a viscoelastic
solid [131, 132].
Many analytical models have been proposed to predict the behavior of
composites based on the analogy between thermal conductivity and diffusivity [133,
134]. The most extensively cited model in polymer/clay nanocomposites is the
Nielsen model, which predicts that relative permeability is only a function of the
aspect ratio at a given loading of clay for all composites [135]. Liu and Wu [136]
have recorded the water absorption curves of PA66 and corresponding
nanocomposites. They concluded that by increasing clay content, the water absorption
at saturation decreases rapidly from 7.6% for PA66 to 5.2% for the nanocomposite
28
containing 5wt.% clay. They suggested that this reduction is due to the presence of
immobilized polymer in the amorphous phase. However, above 5wt.% of clay, the
decrease in the saturation content of water is not obvious, probably because of
aggregation of silicate layers. Also, the diffusion coefficient values show decreasing
trend with increasing clay loading but after 5wt.% of clay, the magnitude of the
decrease is obviously lower.
1.12 Mechanical and wear properties of nanocomposites
Several studies exist in literature wherein polymer nanocomposites with ceramic
nanofillers (both uncoated and coated/functionalized) have been investigated for their
mechanical properties and results have again demonstrated a significant enhancement
in properties. Ceramic nanofillers are mainly considered for enhancing tribological
properties in addition to the regular mechanical characteristics and the fillers in this
category which has been widely explored are aluminum oxide (Al2O3), titanium
dioxide (TiO2) and silicon dioxide (SiO2). Wetzel and co-workers [137] performed a
comprehensive study to analyze the influence of TiO2 and Al2O3 nanoparticles on the
fracture and toughness properties of epoxy and the obtained results were highly
encouraging. It is seen that the inclusion of both Al2O3 and TiO2 nanoparticles into
epoxy resulted in an improvement in the flexural stiffness, flexural strength and
fracture toughness at the same time. The Al2O3 nanoparticles were also found to
additionally improve the fatigue crack propagation resistance in the epoxy
nanocomposites compared to the unfilled epoxy. Similar improvements in the
mechanical performance of polymer nanocomposites due to the influence of ceramic
fillers have been achieved in many other cases too, to mention some of them, (1) the
wear resistance of PTFE (poly tetra fluoro ethylene) improved by over 600 times at
20% filler concentration of s nanoparticles [138], (2) epoxy composites with TiO2
nanoparticles showed a reduction in crack propagation, increased wear resistance
functionality, higher modulus, higher strain to failure and improvements in stiffness
and impact strength [139-141], (3) epoxy-SiO2 nanocomposites displayed improved
friction and wear characteristics as compared to microcomposites at very low filler
loadings and also enhancements in their modulus, microhardness, fracture toughness,
tensile strength, tensile modulus and impact strength [142, 143], (4) the incorporation
of just few volume percent of Al2O3 nanoparticles into epoxy could enhance the
stiffness, impact energy, failure strain, storage modulus, young's modulus, tensile
strength and fracture toughness [144, 145], (5) polypropylene (PP) filled with
nanoSiO2 demonstrated a simultaneous improvement in the modulus, strength and
elongation to break [146].
Elansezhian and co-workers [147] studied the wear and tensile strength
behavior of vinyl ester with silica, alumina and zinc oxide nanofillers and reported
that the addition of silica to the vinyl ester resin significantly improved wear
resistance as compared to other two fillers. The functionalized silica nanoparticles
showed an improved dispersion with vinyl ester resin. Functionalization caused
particle dispersion more uniformly in the polymer matrix. As-received nanoparticles
29
show lower tensile strength whereas functionalized nanoparticles show improved
tensile strength by more than 15% at 5wt.% loading as compared to unfilled resin.
Kinloch’s group [148-151] investigated the effect of nanosilica particles on the
modulus and fracture behavior of epoxy resins and found improvements in the
properties. The maximum improvement observed in the modulus and fracture
toughness was 30% and 140%, respectively, with 20wt.% nanosilica.
The nanosilica/matrix debonding followed by plastic void growth was
identified as the major toughening mechanism [149]. They also used these particles to
successfully develop carbon fiber reinforced epoxy composites with enhanced inter-
laminar fracture properties [151]. Ma and co-workers [152] have reported about 40%
improvement in the tensile modulus and about 130% improvement in the fracture
toughness of epoxy resins with 20wt.% nanosilica. Other authors have also reported
improvements in the modulus and toughness of nanosilica/epoxy nanocomposites
[153-155]. Zhang and co-authors [155] have also observed that significant
improvements in properties were obtained when the interparticle distance was smaller
than the nanoparticle diameter; this was attributed to the three-dimensional network
formed under this condition.
Uddin and Sun developed nanosilica (15wt.%) modified-glass fiber reinforced
epoxy composites and measured the compression, tension, fracture and impact
properties [156, 157]. They improved the longitudinal compression modulus by 20 to
40%, strength by 60 to 80%, longitudinal tensile strength by 11% and transverse
tensile strength by 30% [156]. They also observed improvements in the inter-laminar
fracture and impact properties of the nanomodified-composites [157].
Voigt and co-authors [158] have used colloidal nanosilica particles to develop
nanosilica/epoxy nanocomposite with unique properties in the lithography patterning,
which has applications in the micro- and nano-electromechanical systems. Richard
West and co-workers [159] have developed epoxy-alumina nanocomposites up to
10wt.% through ultrasonic cavitation achieving better dispersion of nanoparticles
which improved the elastic modulus of the polymer nanocomposites and enhanced
stress at 5% strain values.
Cao and co-authors [160] have prepared ternary composites with the
nanoalumina particles dispersed in a binary matrix viz, the modified epoxy resin by
the polyester. They achieved maximum impact strength at 8phr up to 110% more than
that of the binary matrix and 400% relative to that of unmodified epoxy resin.
Similarly, increasing in tensile strength of 44 and 165%, corresponding to those of the
binary matrix and the unmodified epoxy resin was observed. The dielectric loss of the
composites was 10-4
below the temperature of 120°C and they also observed that glass
transition temperature was 119°C.
In recent years [161, 162], polymer is extensively utilized in tribological
applications such as cams, brakes, bearings and gears because of their self-lubricating
properties, lower friction and better wear resistant. The inherent deficiency of
polymers is altered successfully by using various special fillers such as SiC, SiO2,
ZnO, ZrO2, Al2O3 and TiO2 (micro to nanosized particles). More and more polymer
composites are now being used as sliding components which were formerly composed
30
only of metallic materials. Nevertheless, new developments are under way to explore
other fields of application for these materials and to tailor their properties for extreme
load-bearing and environmental temperature conditions [161, 162].
Rong and co-authors [163] have proved that the epoxy/SiO2/TiO2
nanocomposites are effective in lowering the frictional coefficient and wear rate. The
results of these experiments [163] indicate that the wear mechanism of composites
changed from adhesive wear to mild abrasive wear and fatigue wear with increase in
SiO2/TiO2 content. The wear and friction performance of composite have greatly
improved with the addition of SiO2/TiO2 nanoparticles.
1.13 Dynamic mechanical analysis of nanocomposites
Dynamic mechanical properties were evaluated to study the physical, chemical and
structural changes of the polymers and nanocomposites. The glass transition or
secondary transitions yield information on the morphology of polymers were
determined. Viscoelasticity is a characteristic property of polymers and dynamic
mechanical thermal analysis is one of the leading tools for measuring viscoelasticity
of polymers and polymer based composites [164]. Ratna and co-authors [165] studied
dynamic mechanical behaviors of epoxy/clay nanocomposites. They have reported
that incorporation of clay can lead to a promising increase in storage modulus and a
modest increase in Tg of the nanocomposites.
The increase in storage modulus values was explained in terms of the
nanocomposite morphology. Xu and co-authors [166] measured the Tg of the epoxy/o-
MMT nanocomposites using DMA and they observed reduction in Tg with increase in
clay content of epoxy matrix. Chiang and co-authors [167] preformed a series of
experiments on the dynamic mechanical properties (elastic modulus, thermal stability,
and glass temperature Tg of PI (Polyimide)/TiO2 nanocomposites with titania of
different weight percentages and three different PI systems Pyromelliticdianhydride
(PMDA) series; 3, 3´,4,4´-biphenyl tetracarboxylic dianhydride (BPDA) series and
3´,4,4´-benzophenonetetracarboxylic acid dianhydride (BTDA) series. They observed
that TiO2 nanoparticles formed via sol-gel process uniformly disperses in the PI
matrix and result in enhanced dynamic mechanical properties.
For all three series of systems, the elastic modulus increased with increasing
volume fraction of TiO2. For 9% TiO2 additions, the increase in effective elastic
modulus was around 30% over that of the pure PI. The results implied that the
flexibility of hybrid films follows the order: BTDA > BPDA > PMDA. For all the
systems, the storage modulus (E') of hybrids improves with increasing TiO2 levels at
lower and elevated temperatures. But their loss modulus (E") decreased with the
increasing amount of TiO2. Hence, Tg shifted to higher temperatures by increasing the
amount of TiO2 which is much stiffer than the pure PI.
Amit chaterjee and co-authors [168] have reported that with increasing
percentage of TiO2 nanoparticles loading, Tg increases linearly. At 1wt.% loading of
TiO2 nanoparticles (5nm), the epoxy resin gives highest Tg and then starts decreasing.
Therefore by increasing the nanoparticles, the Tg tends to move to higher values
31
relative to the Tg of the pure system (118°C). The nanoparticles obsequiously
influence the Tg and the increase in Tg may be attributed to loss in mobility of chain
segments of the epoxy system resulting from a nanoparticle/matrix interaction.
The improvement of viscoelastic properties was obviously indicated by
incorporation of titania into the PBa matrix [169]. For example, the storage moduli of
the hybrids at room temperature increased with the increase of the titania content and
remained constant up to relatively higher temperature in comparison to the pure PBa
resin. These increments in the storage moduli of the hybrid materials in comparison to
the pure resin indicated that the micro-Brownian motion of PBa segments is restricted
by the homogeneous dispersion of titania into the matrix, leading to reinforced PBa
network. In addition, the Tg of pure PBa (151°C) shifted to 161, 171 and 179°C with
inclusion of 3, 5 and 7wt.% of titania, respectively.
Gefu Ji and co-authors [170] have compared the results of different processing
methods viz., G1 through G5 and reported that G4 has the best storage modulus, loss
modulus and the best mechanical stiffness. Comparing G4 and G5, it was observed
that glass transition temperature of the groups with nanoclay increased by 12°C as
compared to the pure vinyl ester. This is understandable because the elastic properties
of the nanoclay are independent of temperature over temperature range used [170].
1.14 Thermal conductivity and heat distortion temperature
The practical applications point of view the thermal conductivity has been studied
widely. Irwin and co-workers [171] compared the thermal conductivity behaviors of
polyimide filled with micrometer and nanometer sized particles. It was observed that
the thermal conductivity in nanocomposites increased steeply at low filler
concentrations (<5%) as compared to microcomposites. Further, coated nanoparticles
were shown to have a much more pronounced effect on the thermal conductivity as
compared to the nanocomposites (with uncoated nanoparticles) and microcomposites.
In a different study, it was reported by Fan and co-workers [172] that the
addition of aluminum nanoparticles to an epoxy composite which already contained
microparticles of alumina resulted in a reduction in the thermal conductivity of the
final composite material. The authors reasoned that the reduction was due to an
increase in the volume fraction of interfaces in the material due to the presence of
nanofillers. Similarly, a lower thermal conductivity has been recorded for silicone
rubber-SiO2 nanocomposites as compared to microcomposites of the same materials
and at the same filler concentration [173, 174].
In an epoxy system [175] and using silica as the filler (thermal conductivity ~
1.5W/mK), a large volume fraction of fillers (>50%) were required to achieve a small
increase in the thermal conductivity whereas with carbon fibers (thermal conductivity
~800W/mK), the value increased 5 times with respect to the base resin with only
around 30% of filler loading. Similar enhancements in the effective thermal
conductivities of polymer composites were observed in many other systems too, e.g.,
epoxy-alumina [176, 177], polyurethane-alumina [178], polyurethane-carbon fiber
[178], epoxy-AlN [176, 179], PVDF-A1N [179], PVDF-SiC [179], PP-silver [180],
32
epoxy-Al [181], epoxy-CuO (cupric oxide) [182], polyethylene-carbon black [183]
and polyethylene-boron nitride [183].
An important observation which can be made from the thermal conductivity
behaviors in polymer composites is that the increase in the thermal conductivity with
respect to filler concentration is found to be more gradual unlike the existence of a
percolation threshold in the case of electrical conductivity variations [184]. In
addition, a sufficiently high filler loading (at least 50% and above depending on the
filler) is necessary to get a significant increase in the thermal conductivity values.
Apart from the experimental studies, numerous theoretical studies have also been
performed on the effective thermal conductivity of two-phase composites and several
empirical models have been proposed to predict the composite thermal conductivity
[184]. However, these model predictions are not always found to agree with
experimentally obtained results and in a majority of the cases, the agreement has been
observed with one of the models only or at lower filler concentrations depending on
the polymer matrix and the filler material [180-184]. In an interesting result, filler size
has been reported to influence the thermal conductivities of polymer composites arid
for polyurethane-Al2O3 systems, a higher composite thermal conductivity has been
observed with smaller particle sized fillers [178].
Progelhof and co-authors [185] have presented an exhaustive overview on
models and methods for predicting the thermal conductivity of composite systems.
Procter and Solc [186] have used Nielsen model as a predictive tool to investigate
thermal conductivity of several types of polymer composites with different fillers and
have confirmed its applicability.
Nagai and co-workers [187] have observed that Bruggeman model for
Al2O3/epoxy system and a modified form of Bruggeman model for AlN/epoxy system
are both good “prediction theories” for thermal conductivity. Griesinger and co-
workers [188] have reported that the thermal conductivity of low-density poly-
ethylene (LDPE) increases from 0.35W/mK for an isotropic sample, to the value of
50W/mK for a sample with an orientation ratio of 50. The thermal and mechanical
properties of copper powder filled poly-ethylene composites are reported by Tavman
[189] while Sofian and co-workers [190] have investigated the thermal properties
such as thermal conductivity, thermal diffusivity and specific heat of metal (copper,
zinc, iron and bronze) powder filled HDPE composites in the range of filler content of
0-24% by volume. They have reported a moderate increase in thermal conductivity up
to 16% of metal powder filler content.
Tekce and co-authors [191] have reported the strong influence of shape factor
of fillers on thermal conductivity of the nanocomposites, while Kumlutas and Tavman
[192] have reported the results of experimental and numerical studies on thermal
conductivity of particle filled polymer composites. Apart from the experimental
studies, numerous studies have also been performed on the effective thermal
conductivity of two-phase composites and several empirical models have been
proposed to predict the composite thermal conductivity [193].
33
1.15 Co-efficient of thermal expansion
The introduction of well-dispersed inorganic particles into a polymer matrix has been
demonstrated to extremely effective in improving the performance of the polymer
composites. Because of the exceptionally low co-efficient of thermal expansion
(CTE) of silica, which is only 0.5ppm/°C, the silica filled composite materials have
attracted much attention to reduce the CTE of polymer composites and to improve the
mechanical properties. A typical example of silica filled polymer composite in micro
electronics applications is the underfill. Underfill is a layer of adhesive that is applied
between the chip and the substrate to alleviate thermal mechanical stress on the solder
joints in the flip-chip package [194] silica is used as the filler to reduce the CTE for
underfill so as to match the CTE of the solder material to achieve high reliability
[195].
The difference between the thermal expansion co-efficient of a metal and
epoxy insulation in a molded transformer gives raise to partial discharge during
extended use because of voids or interfacial defects on the inner side of an insulating
solid material caused by thermal concentration and expansion [196]. In contrast to
conventional epoxy insulation materials such as aluminum with a thermal
conductivity co-efficient of 2-3ppm/°K, the epoxy insulation shows a value three
times higher (T<Tg).
Liu [197] and co-workers have reported, thermal stability of the phosphorous
containing epoxy resins diglycidylether of bisphenol-A was improved with the
incorporation of the colloidal silica, by blending method with loading as high as
70wt.%. The hybrid materials are cured with commercial curing agents without
altering the curing conditions. The resulting cured epoxy silica hybrid resins (ESHR)
showed good transparency and miscibility as observed with AFM, SEM and TEM. A
depression on the glass transition temperature of the resins was observed, owing to the
plasticizing effect of the colloidal silica. They concluded that, the nanoscale colloidal
silica did not show effectively synergistic effect on char formation and flame
retardance with phosphorus.
Pethrick and co-workers [198] reported on the cure and physical properties of
an epoxy resin created using functionalized nanosilica filler. They showed that, a
decrease in the value of the glass transition temperature (Tg) with increasing silica
level due to cure timings. Dynamic mechanical thermal analysis showed decrease in
the value of the glass transition temperature (Tg) with increasing silica level. They
concluded that, the ability of the nanosilica create a stable network structure by the
variation of the high temperature in the expansion coefficients with increasing silica
level, indicating the effectiveness of the functionalized silica nanoparticles in forming
a network. The network formed during cure in the nanomodified epoxy is unable to
undergo the densification possible in the pure epoxy resin material and explains the
observed lowering of Tg with increasing nanosilica content.
34
1.16 Glass transition temperature analysis using differential scanning
calorimetry
Basara and co-workers [199] have studied the effect of clay type and its content on
the Tg of nanocomposites. They suggested that the Tg of pure epoxy increases from
73°C to 83.5°C with addition of 9wt.% of organically modified MMT (Cloisite 30B)
and to 75°C with 9wt.% of natural clay (Cloisite Na+). They have explained the
results in terms of hindered mobility of polymer chains with addition of clay. Yasmin
and co-workers [200] have observed continuous drop in Tg with increasing clay
content for both nanocomposites and related it to clay aggregates, interface regions
and adhesion problems at the clay-matrix interface. Isik and co-workers [201] have
investigated Tg of layered clay/epoxy nanocomposites. They have reported that Tg
increases with increasing clay content and they have explained this behavior in terms
of restricted mobility of polymer chains due to clay and polymer interaction.
The effect of dodecyl-montmorillonite (DMONT) content on Tg of both rigid-
rod and flexible polyimide film was studied by Magarphan and co-authors [202].
They have suggested that the Tg values of the rigid-rod polyimide are higher than that
of flexible polyimides. Moreover, Tg values of the rigid polyimide nanocomposites
are lower than that of pure polyimides and show variations. The introduction of
vermiculite (VMT) into PP, shows a slight increase in the Tg of the PP matrix to
15.8°C as reported by Tjong and co-authors [203]. The authors have reported that this
supports the fact that in PP/Maleic anhydride (MA)/VMT nanocomposites, the
mobility of the PP chains is restricted by the presence of VMT layers. However, the
melting temperature of the nanocomposites shows variations because the nanoscale
fillers do not alter the crystalline size.
With the introduction of VMT into PVA, a slight increase in the Tg of the
PVA-VMT nanocomposites to approximately 73°C has been reported by Xu and co-
authors [204] and the Tg increased to approximately 3.4°C only for 5wt.% VMT in
PVA. This suggests that the VMT layers are well dispersed in the PVA-VMT
nanocomposites. The amorphous chains of PVA become stable and intercalated
strongly with the VMT layers. In other words, the intercalated VMT can restrict the
motion of the PVA molecular segment. With further increase in clay content, no
further increase in Tg of the nanocomposites was observed. They suggested that this
could be due to the excessive coagulation of VMT in PVA solution and hence it does
not disperse individually to the expected level. Therefore the presence of coagulated
VMT may not be responsible for increase in Tg of PVA.
1.17 Free volume analysis using positron annihilation life time
spectroscopy
Recently, Becker and co-authors [205] have studied the influence of clay on the free
volume properties in cured epoxy and they observed increase in free volume sizes in
the polymer due to the presence of clay. Wang and co-authors [206] have reported
that at low rectorite (MMT) content (0-2.0%), the free volume size in nanocomposites
is nearly the same, but its concentration decreases with increasing content. The
35
exfoliated structure was examined using XRD and interfacial layer formation between
rectorite platelets and epoxy matrix was probed by positrons. However, it is evident
from literature that till date, studies on free volume aspects of polymer-based
nanocomposites are few and are confined to thin polymer films [207, 208] and few
polymers with nanosized spherical fillers [209-211].
Positron annihilation lifetime spectroscopy (PALS) is a versatile and novel
technique which provides information on the nanometer sized free volume cavities
and their concentration from the measured lifetime of ortho-positronium (o-Ps) specie
that predominantly annihilate in the free volume sites of the composite matrix.
Positron lifetime experiments on polymer clay nanocomposites have been reported for
only few polymer clay systems [212-214]. Intercalated polystyrene/clay composites
with a very high clay content (75wt.%) have shown that positron annihilation
behavior in nanocomposites is very similar to the behavior in the clay itself.
Winberg and co-workers [215] have studied the effect of filler content and
filler particle size on the free volume properties and the positron annihilation
characteristics on a series of polydimethylsiloxane (PDMS)/fumed silicon dioxide
(SiO2) composites at temperatures between -185 and 100°C using positron
annihilation lifetime spectroscopy (PALS). The glass transition behavior of the
PDMS/SiO2 composites was determined with differential scanning calorimetry. A
clear influence on the o-Ps lifetime (τ3) in the polymer upon addition of nanosized
fumed SiO2 was observed at all temperatures. A transition in the temperature
dependence of the o-Ps lifetime was observed close to -35°C above which
temperature PDMS exhibits long o-Ps lifetimes. A relationship between τ3 and the
surface tension, equivalent to the behavior of ordinary molecular liquids was observed
in this temperature region. They concluded that, the o-Ps yield was strongly reduced
in the crystallization region and by addition of SiO2. The nonlinear relationship
between filler weight and o-Ps yield could be due to out-diffusion of positrons and/or
o-Ps from the filler particles to the matrix.
Hamdy and co-authors [216] have studied positron annihilation lifetimes
(PAL) for two viscosity-average molecular weights of poly(methyl methacrylate),
PMMA, as a function of temperature. The PAL measurements were performed under
vacuum in the temperature range from 22°C to 150°C with interval of 10°C. The
lifetime spectra were analyzed using two methods: (1) average results of the ortho-
positronium (o- Ps) lifetime and its intensity obtained by PATFIT program and (2) the
o-Ps lifetime and o-Ps hole volume distributions given by Bayes' theorem and the
maximum entropy principle using MELT program. They observed two different
transitions within the temperature range studied. The first is due to the reduction of
non-equilibrium states that are frozen below this temperature. The other is in
agreement with the glass transition temperature of PMMA. They concluded that, the
value of the o-Ps lifetime in the sample with the lower viscosity-average molecular
weight is higher than that with the higher viscosity-average molecular weight while it
increases with increasing temperature. On the other hand, the o-Ps intensity as well as
the relative fractional of the o-Ps hole volume shows behavior in contrast to the o-Ps
lifetime with the viscosity - average molecular weight.
36
Jean and co-workers [217] have investigated surface and interfacial properties
in thin polymeric films using positron annihilation spectroscopy, coupled with a
variable mono-energetic positron beam. They measured free-volume properties from
ortho-Positronium (o-Ps) lifetime and the S parameter of Doppler broadening of
energy spectra from annihilation radiation as a function of the depth and of the
temperature in thin polymeric films. They presented depth profiles of glass transition
temperature and nanoscale layered structures in polystyrene (PS) thin films on the Si
substrate. They observed a significant variation of Tg suppression as a function of
depth in an 80nm polystyrene thin film on Si: 17K lower near the surface and 11K
lower in the interface of the Si substrate than the center of the film or in the bulk.
They concluded that, this depth dependence of Tg suppression is interpreted as a
broadening of free volume distribution in the surface and interfaces.
Wang and co-workers [218] have studied the effects of different dispersion
states of nanolayered OMMT on the positron annihilation parameters and the
mechanical properties for epoxy resin/organic montmorillonite (OMMT)
nanocomposites. They found that the ortho-positronium (o-Ps) intensity decreased
with increasing OMMT content, which indicated that the interaction between the host
and nanofillers restrained the segmental motion, resulting in a decrease of the free
volume. Interestingly, they observed a good correlation between the interfacial
interaction and mechanical properties, suggesting that the dispersion states of OMMT
and interfacial property between clay layers and matrix played an important role in
determining the mechanical properties. They concluded that, he analysis of positron
lifetime results reveals that the dispersion states of nanoscale OMMT layers in epoxy
resin/OMMT nanocomposite play an important role in determining the interfacial
property and the interaction between the OMMT and epoxy matrix. Exfoliated
structure enhances the flexural and impact strengths of nanocomposites due to the
strong interfacial interaction between OMMT and epoxy matrix.
1.18 Fourier transform infrared spectroscopy
Reddy and co-workers [219] have made a comparative study on the structural,
thermal, mechanical and thermo-mechanical properties of ethylene-octene copolymer
(mPE) nanocomposites synthesized with pure nanosilica (NS) and nanosilica-
functionalized with diglycidyl ether of bisphenol-A (ENS) for a loading level of
2.5wt.%. The effects of pure nanosilica (NS) and epoxy resin-functionalized-
nanosilica (ENS) on the structural properties of ethylene-octene copolymer were
analyzed by fourier transform infra red spectroscopy (FTIR), wide-angle-x-ray
diffractometer (WAXD), transmission electron microscope (TEM) and scanning
electron microscope (SEM). Surface functionalization of NS particles with DGEBA
has lead to improved dispersion ENS particle in the mPE matrix supported by FTIR.
This was further supported by TEM study which shows a good dispersion of ENS
particles in case of mPE-ENS. From the FTIR studies, they suggested that, NS is
hydrophilic in nature and the surface of NS particles possess three types of silanol
groups. These are vicinal, geminal and isolated silanol groups (Si-OH). The high bond
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strength of Si-O renders the surface of silica too acidic in nature and as such highly
reactive towards Lewis bases. They reported the reaction between silanol groups of
NS and oxirane group of DGEBA in presence of Lewis acid (SnCl2) in their earlier
communication [220].
Acharya and Prjapathi [221] have studied polymer nanocomposites in the form
of thin films of 30μm by using solution cast method. The samples were prepared with
silver nanoparticles and TiO2 nanopowder and dispersed in polycarbonate (PC). These
composite polymer films were irradiated by various doses of microwaves at 100 to
750W for 10min in commercial microwave oven. The FTIR spectra were taken for
various samples and compared. They concluded that the Ag-PC nanocomposite and
TiO2-PC nanocomposite does not show any formation/deformation of chemical bonds
at microwave irradiation power ranges from 180-750W. The percentage transmission
changes by formation of composites and suggested cluster formation in polymer
films.
Ashok kumar and co-authors [222] have studied room temperature cured
epoxy (LY-556/HY-951) system filled individually with fumed silica (FS) and
modified clay (MC) synthesized by mechanical shear mixing with the addition of tri-
ethylene-tetra-amine (TETA) hardener. They showed that, from the curing studies the
addition of FS in epoxy resin aids the polymerization by catalytic effect and MC
addition does not show any effect in the curing behavior of epoxy polymer.
Thermogravimetric analysis (TGA) shows enhanced thermal stability of epoxy with
FS fillers than that of epoxy with MC fillers. The epoxy with FS fillers shows
considerable improvement in tensile and impact properties over pure epoxy polymer
and epoxy with MC fillers. SEM studies show that addition clay significantly turns
the epoxy system from brittle to ductile nature and this aspect played instrumental
role in scaling up the performance. Epoxy with FS fillers shows enhanced vibration
characteristics than that of the pure epoxy polymer and epoxy with MC fillers and it is
substantiated from reports on FTIR studies that formation of C-H bonds takes place
on the surface of the nanocomposites.
1.19 Objectives of the present research work
The current state of the art of nanodielectric systems has shown promise in terms of
material characteristics which are suitable for many new industrial applications. In
order to fully exploit the opportunities available in “technology development” for
industrial and engineering applications, epoxy matrix was chosen. The nanofillers
which could meet the challenges for industrial applications were silicon dioxide
(SiO2), alumina (Al2O3) and zinc oxide (ZnO). Since behavior of epoxy
nanocomposites with these fillers is not well understood, it was planned to arrive at a
comprehensive model to describe these nanocomposites through well established
evaluation techniques using dielectric, mechanical, wear and thermal properties.
In addition to understanding the material behavior, further insight into material
structure and its correlation to electrical, mechanical, wear and thermal properties was
carried out through studies on interfacial effects of surface modified filler-epoxy
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composites. Also, extensive use of FTIR, DSC and PALS has been made.
1.20 Scope of the present work
The study aims at fabrication of epoxy nanocomposites using three fillers namely
SiO2, Al2O3 and ZnO. Process optimization will be analyzed to establish correlation
between process parameters and performance of nanocomposites.
The epoxy nanocomposites will be thoroughly evaluated through a series of
well planned experiments. The main focus of all the experiments would be to explain
the material behavior in terms of each property studied. Hence the results of series of
experiments carried out, through unique in their own way will be interlinked to
material structure. In this aspect, the approach of the present study is unique and
different from other studies.
The three epoxy nanocomposites namely epoxy-SiO2/Al2O3/ZnO will be
systematically analyzed for dielectric, mechanical, wear and thermal properties.
Additional studies on interfacial phenomenon using PALS, FTIR and DSC etc. would
add value to the understanding of dielectric, wear, mechanical and thermal behavior
of material. Having established material properties an attempt will be made to identify
possible industrial applications for each epoxy nanocomposite.
1.21 Organization of thesis
The thesis has been divided into nine chapters in order to bring out the importance
and significance of various experiments carried out as a part of this investigation.
Since huge experimental data has been accumulated, it is desirable to explain the
observed results in easy and comprehensive manner, keeping the requirements of
“ease of flow” and “clarity of understanding” of the subject matter in mind. Hence
results of dielectrics/electrical, mechanical, tribological, dynamic mechanical and
thermal properties of epoxy based nanocomposites are presented in separate chapters.
Introduction and review of literature on aspects covering epoxy, fillers,
processing methods, interfacing modeling, properties of polymer nanocomposites etc.
are presented in chapter 1. The aim, objectives and scope of the research work is also
discussed under this chapter.
Chapter 2 deals with the details of epoxy and fillers namely silicon dioxide,
alumina and zinc oxide. Further, nanocomposites processing methods adopted for
developing nanocomposites are highlighted and discussed. Details of studies on “as-
cast” surface morphology of the fabricated nanocomposites using Transmission
Electron Microscopy (TEM), X-Ray Diffraction (XRD) and Scanning Electron
Microscopy (SEM) are described and discussed. An overview of the experiments and
techniques used for determination of electrical, mechanical, wear and thermal
properties are also presented in this chapter.
Chapter 3 presents and discusses results of dielectric constant, tanδ,
polarization and depolarization current characteristics of epoxy nanocomposites and
the effects and influence of different experimental parameters on behavior of
nanocomposites are analyzed and discussed.
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The interfacial properties of epoxy nanocomposites are characterized and
presented in Chapter 4. The studies carried out are chemical bonding, glass transition
temperature and free volume measurements. A model is suggested to explain the
observed results based on the dielectric constant and tanδ characteristic of the epoxy
nanocomposites.
Chapter 5 deals with the investigation on dc volume and surface resistivities,
ac breakdown strength, resistance to arcing, tracking and tracking index
characteristics of epoxy nanocomposites. The influence of nanofiller loading, filler
size are analyzed. The correlation between breakdown strength and tracking index of
the nanocomposites is carried out with the existing models. The effect of contaminant
flow rate and electric field, variation on leakage current is investigated by employing
dimensional analysis technique.
In Chapter 6, the effect of surface treatment of nanoparticles and the influence
of water absorption on epoxy nanocomposites are examined. Results of contact angle
measurements, glass transition temperature, dielectric constant and tanδ, ac
breakdown strength and results of resistivity measurement are presented and
discussed in this chapter. In addition, the influence of water on the dielectric
properties of epoxy nanocomposites explained by water shell model.
Experiments were carried out to understand the mechanical and wear behavior
of the epoxy nanocomposites and a correlation is attempted between electrical and
mechanical properties. These results are presented and discussed under Chapter 7.
Chapter 8 focuses on the influence of nanofiller loading on dynamic
mechanical and thermal properties of the epoxy nanocomposites. The influence of
addition of nanofiller on storage modulus and heat deflection temperature are
analyzed and correlated with suggested models.
Chapter 9 summarizes the conclusions of the present research work and
highlights the scope for future research work in the area of epoxy nanocomposites.
Some of the important industrial applications of the epoxy nanocomposites are
suggested and discussed in this chapter.
The study has helped in understanding fabrication and evaluation methods for
nanocomposites, in addition to identifying the gaps in technology development. With
this background, details of methods of fabrication and experimental methods are
discussed and presented in Chapter 2.