Mechanical and dielectric properties of epoxy-clay nanocomposites
A. Guevara-Morales1,2* and A.C. Taylor1
1. Department of Mechanical Engineering, Imperial College London, South Kensington Campus, London
SW7 2AZ, UK.
2. Currently at Department of Mechatronics Engineering, ITESM-CEM, Atizapán de Zaragoza, Estado de
México 52926, México.
* Corresponding author. A. Guevara-Morales. Tel.: +52 5558645555. Email: a.guevaram@itesm.mx
A.C. Taylor email: a.c.taylor@imperial.ac.uk
Mechanical and dielectric properties of epoxy-clay nanocomposites
Abstract
Epoxy-clay nanocomposites were prepared using two types of surface-treated montmorillonite (Closite 30B
and Nanomer I28E). Wide angle X-ray scattering showed that all the nanocomposites had an intercalated
structure. Improvements in tensile and fracture properties were found. The pure epoxy polymer was very brittle
with a fracture energy, Gc, of 131 Jm-2. The addition of the nanoclays significantly increased the value of Gc,
up to 240 Jm-2 for 5wt% C30B. The toughening mechanisms acting in the nanocomposites were identified
using scanning electron microscopy as crack deflection and plastic deformation of the epoxy matrix around
the clay platelets following debonding. From electrical testing, the permittivity and loss angle of the
nanocomposites decreased and their breakdown strength increased as desired for insulation applications. The
breakdown strength of the pure epoxy was found to be 11.7 kVmm-1, while for a 2 wt% C30B nanocomposite
it increased to 14.7 kVmm-1. It was concluded that the restriction of chain mobility inhibited electrical
polarisation and thus decreased the permittivity and loss angle. The electrical damage zone was analysed using
scanning electron microscopy. It was found that the higher resistance to surface degradation by partial
discharges and the creation of a tortuous electrical path, which delayed the propagation of the electrical tree,
were the main factors which improved the breakdown strength of the nanocomposites.
Keywords: Epoxy-clay nanocomposites; montmorillonite; toughening; dielectric properties; breakdown
strength; electrical tree; partial discharges.
1. Introduction
Thermosetting epoxy resins are widely used for many applications, including as high strength adhesives,
surface coatings, durable laminates and lightweight foams [1]. They are also indispensable materials for
insulation systems in power equipment and heavy apparatus such as dry type transformers, rotating electrical
machines and switchgear [2]. During the operation of such power equipment, the insulator is subjected not
only to electrical stress, but must withstand thermal and mechanical stresses as well as tribological loads [3].
In recent years, polymer nanocomposites reinforced with surface-treated montmorillonite nanoclays have
received significant attention due to the substantial improvements they can achieve in mechanical, thermal and
barrier properties when compared to the unmodified polymer [4-6]. Furthermore, studies have reported [7,8]
that the addition of small amounts of nanofillers can improve the resistance to partial discharges and electrical
treeing growth of the polymer matrix. The selection of layered silicate nanoclays as reinforcements is highly
attractive because of their relatively low cost, high thermal inertness, ready availability and environmentally
friendly characteristics [9].
Significant decreases in permittivity and loss angle have been reported [10,11] with the addition of small
amounts of nanoparticles (~5wt%), which are attributed to the restriction of polymer chain mobility by the
nanoparticles. This restriction inhibits electrical polarisation because the material loses its freedom to relax
under an applied voltage, and thus the permittivity and loss angle decrease. It has also been reported [11] that
at higher clay contents the values of permittivity increase. It has been suggested [11] that this is due to the
presence of accidental inclusions, imperfections or inhomogeneities in the nanocomposite that are technically
difficult to avoid.
It has also been reported [12] that epoxy-clay nanocomposites show better AC and DC breakdown strength.
These improvements are usually related to the electrical tree growth behaviour in the nanocomposites. It has
been suggested [9] that the nanoparticles act as scattering sites that decelerate the electrons and carriers and
hence increase the breakdown voltage. Other authors [11] suggest that the presence of nanoparticles changes
the charge distribution of the material, acting as trapping sites and decelerating the propagation rate of the
electrical tree.
This study investigates the effect of the addition of commercially available nanoclays has on the mechanical,
fracture and dielectric properties of an epoxy polymer. Scanning electron microscopy is used to investigate the
toughening and failure mechanisms.
2. Experimental
2.1 Materials
The epoxy used was a digylcidylether of bisphenol A (DGEBA) (Araldite MY750, Huntsman), with an
anhydride hardener (Aradur HY917, Huntsman) and imidazole accelerator (DY070, Huntsman), which were
mixed in proportions of 100:90:1 by weight. Two organo-modified montmorillonite (MMT) clays were used,
as shown in Table 1; the octadecyltrimethylammonium-treated Nanomer I28E supplied by Nanocor, and the
methyl, tallow, bis-2-hydroxyethyl, quaternary ammonium-treated Cloisite 30B supplied by Southern Clay
Products. These were selected by performing compatibility tests on a range of commercially-available clays,
which showed the relative stability of the dispersion of the clays in the DGEBA.
Table 1. Organo-modified montmorillonite clays used.
Clay
Particle size
Supplier
Modifier
Cloisite 30B
<10µm
Southern Clay Products
MT2EtOH: methyl, tallow, bis-2-hydroxyethyl,
quaternary ammonium;
Nanomer I28E
15-20 µm
Nanocor
ODTMA: Octadecyltrimethylammonium
2.2 Preparation of Nanocomposites
Nanocomposites were prepared by incorporating 1-5 wt% of clay into the epoxy system described above. The
required amount of epoxy resin was poured into a beaker and heated to 60 °C to reduce its viscosity. The dried
clay was sieved using a 300 m mesh to remove any agglomerates, and gently added to the warm epoxy. The
mixture was vigorously hand-stirred using a spatula for a few minutes. An opaque, foamy and viscous mixture
was formed. The mixture was then degassed in a vacuum oven for 10 minutes at 60 °C. The hardener and
accelerator were added, and the mixture was again hand-stirred with the spatula. The mixture was then placed
in the vacuum oven at 60 °C for 1 hour to degas, producing a final transparent solution. The degassed mixture
was poured into a steel mould coated with release agent (Frekote 700NC, Henkel), and cured at for 2 hours at
60 °C, followed by 10 hours at 110 °C. The mould was left in the oven to cool slowly to room temperature,
and the cured plate removed.
2.3 Characterisation
Wide angle X-ray scattering (WAXS) was performed using a Philips PW1700 automated powder
diffractometer equipped with CuK radiation (wavelength 1.54 Å) operating at 40 kV and 40 mA with a
secondary graphite crystal monochromater. Diffraction patterns were collected between 3° and 25° with a step
size of 0.04° and scanning speed of 0.48°/min. Clay samples were prepared by mounting and pressing the clays
into an aluminium holder with a glass back support. Nanocomposite samples were prepared in moulds of 50 x
45 x 2 mm for X-ray scattering.
The glass transition temperature, Tg, was measured using a Triton TT Dynamic Mechanical Analyser. Moulded
prismatic samples of 45 x 3 x 2 mm were loaded under dual cantilever bending mode. Scans were performed
over a temperature range from 40 to 180 °C under controlled sinusoidal strain, at a frequency of 1 Hz and a
heating rate of 2.5 °C/min. The Tg was determined from the position of the peak of the tan  curve, and the
typical standard error was ± 1%.
2.4 Mechanical Testing
Tensile dumbbell specimens were machined from the moulded plates and tested at room temperature according
to the BS ISO 527-2 standard (test specimen 1A). The tests were performed at a constant displacement rate of
1 mm/min. Three replicate samples per formulation were used. The elastic modulus, E, and tensile strength
were calculated.
Fracture specimens were machined from moulded samples and tested according to the ISO 13586 standard at
room temperature of dimensions 70 x 15 x 5 mm. Single-edge notch bend (SENB) tests were performed at a
constant displacement rate of 1 mm/min. An initial crack was created by tapping a razor blade into the
machined notch tip to generate a natural crack which was approximately half the height of the specimens
before testing. Six replicate samples per formulation were used. The fracture toughness, Kc, was calculated
and the fracture energy, Gc, was obtained from Kc according to the standard.
2.5 Electrical Testing
Electrical tests were performed according to the IEC 60243-1 (1998) standard. Values of permittivity, loss
angle, resistivity and breakdown strength were obtained. A plate-plate electrode geometry was used. Tests
were carried out at a power frequency of 50 Hz, and at ambient conditions (19 °C; 55% R.H.).
2.6 Scanning Electron Microscopy
A JEOL JSM-5300 scanning electron microscope was used to examine the fracture surfaces and to characterise
the electrical trees. The samples were sputter-coated with a thin layer of gold before observation to eliminate
charging. For the visualization of the electrical trees, a technique proposed by Vouyovitch [13] was used. This
consists of making a straight saw cut on each side of the damage zone, purposely avoiding the damage, and
then bending the plate at room temperature to break it so that the crack propagates through the damage zone.
3. Results and Discussion
3.1 X-Ray Diffraction
The (001) basal reflections measured by WAXS were used to obtain the interlayer spacing of the clays. The
results are summarised in Table 2. The (002) basal reflections of the 5 wt% nanocomposites were used to
obtain the interlayer spacing of the clays in the epoxy matrix. The nanocomposites showed a further increase
in interlayer spacing, as expected due to the intercalation of the epoxy between the clay platelets [14]. Both
nanocomposites containing 5 wt% of organo-modified montmorillonite showed similar d-spacings.
Table 2. Interlayer spacings of clays and nanocomposites, and glass transition temperatures of epoxy and
nanocomposites.
As-received clay
5 wt% Nanocomposite
Tg
d-spacing (Å)
d-spacing (Å)
[°C]
Epoxy
N/A
N/A
139.3
C30B
17.7
32.5
139.8
I28E
22.6
34.5
145.0
Sample
3.2 Dynamic Mechanical Analysis
The glass transition temperature, Tg was determined from the peak of the tan  curve at a frequency of 1 Hz,
and the results are presented in Table 2. A Tg of 139.3 °C was measured for the unmodified epoxy, values of
139.8 °C and 145.0 °C were found for the C30B and I28E nanocomposites respectively. The glass transition
temperature can vary with various factors such as the molecular weight, crosslink density, free volume, and
the properties of the interface layer [6]. Some authors [9] have reported an increase in Tg with the addition of
clays, which is attributed to the good adhesion of the polymer matrix and the particles acting as restrictions for
chain mobility. Others [14,15] have reported an initial increase in Tg up to a certain amount of clay when Tg
starts decreasing. This decrease in Tg was related to a change in chain kinetics in the regions surrounding the
nanoparticles, which leads to a lower crosslink density [6].
It is observed that Tg increases with increasing d-spacing, as the Tg for the 5 wt% I28E nanocomposite (dspacing: 22.6 Å) is significantly higher than that for the C30B (d-spacing: 17.7 Å), see Table 2. This is an
indication that the clay platelets may be acting as a restriction to polymer chain mobility, and that the more
intercalated or exfoliated the clays are, the more restriction they generate. The increase in Tg for both
nanocomposites indicates that there is good adhesion between both clays and the epoxy matrix.
3.3 Mechanical Properties
Tensile tests results are summarised in Table 3. It was found that the modulus of the epoxy resin (3.53 GPa)
increased slightly when nanocomposites were formed (3.58-3.66 GPa). The exception was the 5 wt% I28E
nanocomposite, where the modulus was decreased, which can be attributed to voids or defects in the material.
The standard deviations of these data vary from 2 to 6 % of the mean value.
It has been shown that the elastic modulus of a polymer is greatly improved when nanocomposites are formed
based on layered silicates [16]. This improvement in elastic modulus is attributed mostly to the much higher
stiffness of the clay particles, but also to the restriction the mobility of the polymer chains under load and by
the clay platelets [17-19]. Orientation of the clay platelets along the loading direction will also contribute to
the increase in elastic modulus [19]. However, Kinloch and Taylor have shown that the orientation of the
platelets in such cast plates is not strong [20], and so a random orientation can be assumed.
Using the Halpin-Tsai model [20], considering a random orientation of the clay platelets, an elastic modulus
of 3.62 GPa was predicted, which agrees well with the measured values. For the calculations a clay modulus
of 172 GPa [20] and an aspect ratio of 15 [21] were used.
Table 3. Elastic modulus, tensile strength, fracture toughness and fracture energy of epoxy and its
nanocomposites. Mean and standard deviation (SD) of the data are shown.
Sample
Epoxy
2wt% C30B
5wt% C30B
5wt% I28E
E [GPa]
Mean
SD
3.53 E 0.14
3.58
0.03
3.66
0.03
3.37
0.09
Tensile Strength [MPa]
Mean
SD
Tensile strength
46.5
6.1
54.2
5.6
54.3
8.3
50.3
7.4
Kc [MPam1/2]
Mean
SD
0.68 Kc 0.05
0.92
0.95
0.83
0.08
0.05
0.04
Gc [Jm-2]
Mean
131
207
240
179
Similar results as for the elastic modulus were found for the tensile strength. As shown in Table 3, all the
nanocomposites have a higher tensile strength (50.3-54.3 MPa) than the epoxy matrix (46.5 MPa). Note that
for brittle materials such as these, the tensile strength is strongly dependent upon the surface finish of the
specimens [20], and that the standard deviations of these data were approximately 10-15% of the mean value.
The load versus displacement behaviour was typically linear elastic, with no evidence of yielding.
The fracture toughness, Kc, of the epoxy and its nanocomposites were obtained from the SENB tests. A fracture
toughness of 0.68 MPam1/2 was measured for the epoxy, which is typical for such a crosslinked polymer [22];
and an increase was found for all its nanocomposites, with fracture toughness from 0.83 to 0.92 MPam1/2, as
shown in Table 3. The epoxy resin has a fracture energy, Gc, of 131 Jm-2, which is typical of a brittle
thermosetting polymer [23]. An increase in Gc was found for all the nanocomposites. The addition of 5 wt%
of C30B almost doubled the fracture energy of the unmodified epoxy to 240 Jm-2. It has been reported that the
addition of nanoclays has a toughening effect on the epoxy resin, which is attributed to crack deflection and
plastic deformation around the nanoparticles [20,24]. Scanning electron microscopy (SEM) was used to
confirm whether these mechanisms were present for the nanocomposites, and the obtained micrographs are
shown in Fig. 1.
The fracture surface of the epoxy is shown in Fig. 1a. It is a smooth and glassy surface, which is typical of the
fracture surfaces of crosslinked thermosets, and shows that little plastic energy dissipation accompanied the
fracture process [25]. The fracture surfaces of the nanocomposites are shown in Fig. 1b-d. These surfaces are
rougher than the epoxy matrix, and as the clay content increases, the roughness of the surface also increases,
which is in agreement with the increased Gc values reported in Table 3, as rougher fracture surfaces indicate
more plastic energy dissipation.
One possible toughening mechanism is explained by the crack deflection model proposed by Faber and Evans
[26], which assumes that as the crack approaches a particle in a material it is deflected by the particle. Crack
deflection will increase the roughness of the surfaces through tilting and twisting of the crack front, increasing
the fracture surface area and thus increasing the fracture energy of the material. In the crack deflection process
modes I and II occur, and in the crack twisting process modes I and III are present. From fracture mechanics
it is known that more energy is required to cause fracture in modes II and III than in mode I [21]. It is important
to mention that the crack deflection theory cannot fully explain the toughening effect in the nanocomposites
unless the particles are larger than the crack opening displacement [26], which is the case for these
nanocomposites [27].
It can also be seen from Fig. 1 that the roughness of the surfaces depends on the type of clay, and that the 5wt%
C30B nanocomposite has a rougher surface when compared to the 5 wt% I28E nanocomposite, which is also
in agreement with the obtained Gc values. Faber and Evans [26] also concluded that the aspect ratio of the
particles has a significant effect on the toughening mechanisms of particulate composites, which may explain
the difference in fracture energy for the C30B and I28E nanocomposites. Note that according to the
manufacturers’ data (Table 1), the particle size of the C30B clay (10 µm) is smaller than the I28E clay (15-20
µm), however, the I28E interlayer spacing is greater (Table 2), which will reduce the aspect effective ratio of
the intercalated particles.
The other toughening mechanism is debonding followed by plastic deformation of the epoxy around the
nanoparticles (i.e. void growth) [20]. Debonding of the clay platelets from the epoxy leads to the formation of
holes. These holes then grow plastically due to the triaxial stress state in the plastic zone at the crack tip,
absorbing energy and increasing the toughness. In Fig. 1b-d, these holes are visible as black regions. Thus, it
is concluded that the plastic deformation of the epoxy matrix around the debonded particles is another
toughening mechanism in these nanocomposites.
3.4 Electrical Properties
The permittivity, loss angle and resistivity values of the epoxy and its nanocomposites are presented in Table
4. The permittivity and loss angle measurements of the 5 wt% I28E nanocomposite were very unstable and
reliable values could not be recorded.
Table 4. Dielectric properties and breakdown strength of epoxy and its nanocomposites.
Permittivity
Resistivity [.m]x1014
Mean
SD
Mean
SD
Epoxy
1.87
0.57
5.30
0.57
1wt% C30B
1.87
0.13
5.25
0.07
2wt% C30B
2.83
1.36
5.35
0.07
5wt% C30B
7.13
2.28
5.27
0.18
5wt% I28E
7.29
0.51
N/D
SD: Standard Deviation; N/D: Not determined
Sample
Loss angle
Mean
SD
0.00371 0.0001
0.00295 0.0002
0.00309 0.0013
0.00471 0.0031
N/D
-
Breakdown Strength [kV/mm]
Mean
SD
11.7
0.8
13.4
2.5
14.7
0.5
13.0
0.3
13.5
1.3
Although inorganic fillers have a higher permittivity values than the polymer matrix, in the range of 5 to 8
[10,28] compared with the 5.30 value for the epoxy (Table 4), the permittivity usually decreases in
nanocomposites when compared to the unmodified polymer [9,11,29]. This reduction is related to the
restriction of polymer chain movement by the nanoparticles. Here the material loses its freedom to relax under
an applied voltage, and the polymer chains are not able to contribute to the electrical polarization [30].
However, the higher values of permittivity which were measured, such as for the 2 wt% C30B, are related to
inhomogeneities or defects in the nanocomposites (e.g. voids, air bubbles) or agglomeration of the
nanoparticles in the polymer matrix [11]. The loss angle also decreased in the nanocomposites. The loss angle
is sensitive to the existence of mobile ionic impurities and polar radicals with dipole moments [11], and its
reduction is related to the restriction in polymer chain motion, like the decrease in permittivity.
From the electrical tests it was found that the breakdown strength of all the nanocomposites was higher than
that of the epoxy matrix. The breakdown strength is defined as the maximum electric field strength that an
insulating material can withstand intrinsically without breaking down [31]. The term breakdown is used to
describe processes in which a considerable current increase results from a small voltage change, and in the
case of polymeric insulators it is always catastrophic, irreversible and destructive, creating a breakdown
channel between the electrodes [7]. This increase in breakdown strength is attributed to a broadening of the
electrical tree and a decrease in the tree propagation rate that may arise from the resistance to partial discharges,
which will be discussed in the following section.
3.5 Resistance to Partial Discharges
To explain the increases in performance, it is necessary to consider the mechanisms due to the nanoclays within
the breakdown process. It has been found that the addition of organo-modified clays improves the resistance
to surface degradation by partial discharges [32]. This improvement can be explained with the multi-core
model proposed by Tanaka [28], which is based on the proposal that the interface is divided in three layers:
the bonded layer, the bound layer and the loose layer (Fig. 2); and that an electric double layer is superimposed
on all of them. The bonded layer is in contact with the particles and it consists of immobile polymer chains
bonded to the particle surface. The bound layer is adjacent to the bound layer and consists of less ordered
polymer chains. Finally, there is the loose layer, which is considered to be an amorphous region with polymer
chains which are relatively free to move.
The model suggests that the electrical tree is attracted to the particle due to its higher permittivity when
compared to the epoxy matrix. When the electrical tree reaches the loose layer of the interface it propagates
through it. However, when it reaches the stronger bound layer it cannot degrade this bound layer so easily and
so the electrical tree prefers to crawl along the loose layer until it reaches the epoxy matrix again. Due to the
close packing of the particles and the small interfiller distance, the electrical tree is rapidly attracted by another
particle, repeating the process and creating a tortuous or zigzag path that will increase the breakdown strength
(see Fig. 2). Thus the role of the nanoclays is not exactly as a barrier, as they attract the partial discharge due
to their higher permittivity, but due to the higher partial discharge resistance the partial discharge prefers to
crawl along the interface, increasing the length of the tree and delaying its propagation.
3.6 Electrical Treeing
Electrical treeing is an electrical pre-breakdown phenomenon in which fine erosion channels propagate through
the insulator material, in a path resembling tree branches, due to partial discharges. The process can be divided
into three stages: inception, propagation, and runaway or final breakdown, which are explained in detail by
Dissado and Forthergill [7]. The inception of electrical trees involves bond breaking and relies on electrons or
carriers whose kinetic energies are sufficient to cause bond scission (3 to 4 eV), but lower than that required
for ionization (8 eV), that can cause electrical breakdown. According to the partial discharge mechanism, the
main step in tree inception is the establishment of a path consisting of voids or defects aligned between the
electrodes. After the inception of the tree (the formation of the first branch), the electrical tree propagation
exhibits a decelerating growth rate that is associated with the branch tip splitting (branching of the tree),
extinction of channel discharges and the lack of a vent to release the gases produced by the discharges. Note
that breakdown is not always an instantaneous process after the first branch reaches the opposite electrode; it
usually requires that the main branch or channel reaches a certain diameter. During the creation of the first
branch, electrons or carriers are extracted from the matrix, leading to the formation of positively charged bands
along the channel length. When these positive bands reach a certain charge density, they produce a local field
which is sufficiently high to attract and neutralise the electrons. At this stage there is not forward erosion, but
a lateral one that raises the possibility of branching and widening of the channel. At the runaway stage the
main channel becomes so wide that charge repulsion will force the cations in the walls to run through it acting
as a conducting projection, higher kinetic energies will be reached, causing the irreversible failure of the
insulator. It has been found that cracking is almost instantaneous once large discharges occur. The faster the
inception and the propagation, the faster the final breakdown will be, and hence the lower the breakdown
strength of the material.
Scanning electron microscopy (SEM) was used to examine the electrical trees in the tested samples, and images
are presented in Fig. 3. The electrical tree branches are clearly observed in the nanocomposites, extending all
over the damage zone (Fig. 3b-e). In the case of the epoxy sample (Fig. 3a), the main channel, indicated by the
black arrow, occupies a great part of the damage zone with a diameter of approximately 400 m at the top and
1000 m at its centre. The diameter of the main channel in the nanocomposites is smaller, being approximately
250 m in all cases.
It is important to notice that the electrical trees in these samples do not have the common “tree” shape with
branches radiating from a point towards the opposite side (see schematic in Fig. 3). Some of the branches are
oriented towards the top surface while others are oriented towards the bottom surface. This is an effect of the
electrode configuration. A needle-plate electrode configuration is commonly used when the electrical tree is
to be analysed. When the applied AC voltage generates a current from the top to the bottom surface, it
propagates from the tip of the needle to the opposite plate. During the other half of the cycle, the current is
reversed, but the path will be the same because the electrons are attracted by the tip of the needle. This generates
the classic tree shape. When a plate-plate electrode configuration is used, as in the present work, the electrons
are not attracted by a single point (Fig. 3c-d) but to a larger area, and therefore the tree propagates towards
both sides. It is important to note that with the needle-plate electrode configuration the tree will be initiated at
the tip of the needle, whereas in the case of plate-plate electrode configuration it will initiate in the vicinity of
the major flaw [33], and hence plate-plate electrodes give a more realistic value for the breakdown strength.
From Table 4 it can be seen that the addition of the organo-modified clays improved the breakdown strength
of the epoxy, which may indicate that both the inception and propagation stages of treeing were delayed with
the addition of the clays. The inception stage is determined by the yield strength of the material and by the free
volume [7]. According to the results of tensile strength presented in Table 3, all the nanocomposites have a
higher tensile strength than the epoxy, indicating that the inception stage may be delayed and hence the
breakdown strength of the material is increased. As the behaviour of the epoxy resin and its nanocomposites
is purely linear elastic, the tensile strength was considered as equivalent of the yield strength. Although the
free volume in the epoxy and its nanocomposites were not analysed in this study, the literature reports an
increase of between 2 and 3 % in the free volume for nanocomposites [11,34], which may indicate a decrease
in breakdown strength due to the larger space where the electrons will be able to move and acquire energy for
bond scission. However, the increase in Tg with the addition of nanoclays (Section 3.2) may be considered as
a reduction in free volume, as they act as a restriction to the mobility of the electrons. An important
consideration here is that the increase in free volume is mainly localised at the interface of the nanoparticles
[28] and when the accelerated electrons or carriers move through the interface they can be decelerated via
collision with the inner layers of the interface, resulting in an increase in breakdown strength. The electrons
will not be able to accelerate again due to the small interparticle distance.
As mentioned above, during the propagation stage the positively charged main channel attracts the electrons
or carriers to its walls, producing a lateral erosion that increases the possibility of branching and widening of
the channel [7]. In the case of the epoxy, this lateral erosion increases the diameter of the main channel, as
observed in Fig. 3, due to the material’s low resistance to surface degradation caused by partial discharges.
When the main channel becomes so wide, charge repulsion will force the cations in the walls to run through it
acting as a conduction projection, reaching kinetic energies high enough to cause bond scission, degrade the
polymer wall and cause the breakdown of the material.
Figure 4 shows the inner surfaces of the electrical tree main channel for the epoxy and its 5 wt% I28E
nanocomposite. Note the significant difference between the rough channel surface for the nanocomposite and
the smooth surface for the epoxy sample. This can be explained by the surface degradation process occurring
in the epoxy matrix and its nanocomposites. Erosion due to partial discharges is considered to occur at the
same time in both specimens (pure epoxy and its nanocomposite), but in the nanocomposites the nanoparticles
offer more resistance to partial discharges, as explained above. Therefore, the nanoparticles are exposed on the
surface, which gives the rougher and whiter appearance in the SEM images. This layer of exposed
nanoparticles will impede further erosion of the surface due to the higher erosion resistance of the clay platelets
[32] and decelerate the widening of the tree channels [8], slowing down the speed of widening of the channels
that lead to the final breakdown.
Some cracks in the walls of the main channel are observed in Fig. 4a. It is thought that due to the lower fracture
toughness of the epoxy resin (see Table 3) the lateral impacts of the electrons in the walls of the epoxy matrix
initiated cracks in the surface of the channel. These cracks then acted as free volume sites [7] and allowed the
branching of the tree. Note that the breakdown strength is not determined by the number of branches, but by
the ease with which those branches propagate. In the case of the epoxy, the low resistance to surface
degradation will facilitate the propagation of the tree.
In Fig. 5a the electrical tree branches of the 1 wt% C30B nanocomposite are shown at a higher magnification.
Although these branches may be seen as an effect of the fracture surface rather than branches of an electrical
tree, the damage zone (where the branches are) is actually part of a crater or depression, as can be observed in
Fig. 3c. Therefore, the branches have not been affected by the way the samples were prepared for SEM
observation, as cutting the damage zone was carefully avoided (Section 2.6). In Fig. 5b a depression is clearly
observed, which is the boundary between the crater of the electrical damage zone and the fracture surface
created when the samples were prepared for microscopy.
It has been suggested that other factors such as the state of stress and the thermal properties of the
nanocomposites can contribute to the enhancement of breakdown strength [33,35]. Due to the smaller thermal
expansion coefficient and lower linear shrinkage ratio of the nanocomposites when compared to the
unmodified epoxy, their state of stress is lower as shown by Imai et al. [35]. This lower state of stress may also
contribute to the enhancement of the breakdown strength [7]. In the case of the thermal properties, it has been
reported that the addition of organo-modified clays allows a better dissipation of the heat generated by the
applied electric field [33], decreasing the possibility of thermal breakdown.
In future, a more detailed investigation of the damage zone is required, especially at the inception and
propagation stages (i.e. when the final breakdown has not occurred yet). This can be done by testing the
specimens at different voltages and for different times in order to be able to determine a sequence of failure
mechanisms that are present in the nanocomposite electrical degradation. The samples presented in this study
were tested until the final breakdown occurred, and therefore the damage zones shown in Figs. 3-5 present
multiple failure mechanisms such as electrical breakdown and fracture. This multiple-failure damage zone
complicates the visualisation of the electrical tree. If the main interest is on the electrical treeing process, it is
recommended to change the configuration of the electrical tests from a plate-plate electrode configuration to a
needle-plate configuration. This will produce a divergent field instead of a quasi-homogeneous one and the
tree will grow more clearly.
4. Conclusions
Epoxy-clay nanocomposites were prepared using two organo-modified montmorillonites. The elastic modulus
of the nanocomposites was higher than that of the epoxy matrix, due to the much higher stiffness of the clay
when compared to that of the epoxy, and to the restriction of polymer chain mobility due to the presence of
the clay. Results showed that the fracture energy of the epoxy (Gc = 131 Jm-2) was increased for the
nanocomposites (to 180-240 Jm-2). The toughening mechanisms acting in the nanocomposites were identified
as crack deflection and debonding followed by plastic deformation around the clay platelets.
The permittivity and loss angle of the epoxy, measured as 5.30 and 0.0037 respectively, decreased in the
nanocomposites, reaching 5.25 and 0.0029 for the 1 wt% C30B nanocomposite. As these two properties are
related to the electrical polarisation, it is likely that the restriction of chain mobility inhibited this polarisation
and thus decreased the permittivity and loss angle as desired for insulator materials. An increased breakdown
strength for all the nanocomposites was found, with the epoxy having a breakdown strength of 11.7 kVmm-1,
and its nanocomposites having values from 13.0 to 14.7 kVmm-1. This improvement was due to the
nanocomposites’ higher resistance to surface degradation by partial discharges, which delayed the propagation
of the electrical tree. The clay particles attract the electrical tree due to their higher permittivity, but due to
their higher degradation resistance the electrical tree prefers to crawl along the weaker layer of the interface
(loose layer) until it reaches the epoxy matrix that is easier to erode. As the interparticle distance in
nanocomposites is very small, the electrical field is attracted by another particle very soon and the process is
repeated. This zigzag path increases the length of the tree and delays its propagation.
Acknowledgements
Dr. Guevara-Morales gratefully acknowledges the financial support from the Mexican National Council for
Science and Technology (CONACyT). Some of the equipment used in this work was provided by Dr. Taylor’s
Royal Society Mercer Junior Award for Innovation.
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Figures
Fig. 1 SEM images of the fracture surfaces of (a) epoxy, (b) 2wt% C30B, (c) 5wt% C30B and (d) 5wt% I28E
Fig. 2 Partial discharges (PD) degradation path in nanocomposites, drawn based on [28]
Fig. 3 SEM images of the electrical trees in (a) epoxy, (b) 5 wt% I28E, (c) 1 wt% C30B, (d) 2 wt% C30B and
(e) 5 wt% C30B. (Arrow indicates the main channel or branch of the electrical tree)
Fig. 4 SEM images of the electrical tree main channel in the (a) epoxy and (b) 5wt% I28E samples
Fig. 5 Electrical damage zone in the 1 wt% C30B nanocomposite, (a) electrical tree branches and (b) crater