1

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,

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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.

15

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.