Functional nanomaterials for electric power industry

661
Functional nanomaterials for
electric power industry
Working Group
D1.40
August 2016

FUNCTIONAL NANOMATERIALS
FOR ELECTRIC POWER INDUSTRY
WG D1.40
Members
M. Fréchette, Convenor (CA), A. Vaughan, Secretary (GB),
A. Allais (FR), U. Häring (DE), T. Andritsch (GB), I. Bergmann (AU), J. Castellon (FR),
A. Cristiano (FR), E. David (CA), M. Darques (FR), V. Englund (SE), D. Fabiani (IT),
A. Germano (BR), S.J. Han (US), F. Perrot (GB), N. Quirke (GB), C. Reed (US), T. Shimizu (JP),
S. Sutton (GB), T. Tanaka (JP), J. Weidner (DE)
Copyright © 2016
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ISBN: 978-2-85873-364-4

FUNCTIONAL NANOMATERIALS FOR ELECTRIC POWER
Page 2
FUNCTIONAL
NANOMATERIALS FOR
ELECTRIC POWER INDUSTRY
Table of Contents
EXECUTIVE SUMMARY...............................................................................................................................3
1 Introduction..........................................................................................................................................3
2 Some empirical guidelines to functionality ..........................................................................................5
2.1 Basic Considerations..........................................................................................................................5
2.2 Fillers..................................................................................................................................................8
2.3 Guiding Considerations....................................................................................................................12
2.4 Closing Remarks ..............................................................................................................................23
2.5 References .......................................................................................................................................23
3 Strategies towards novel nanostructured functional dielectrics........................................................28
3.1 Introduction.......................................................................................................................................28
3.2 Polymer Blends – Functionality through Mixing ...............................................................................29
3.2.1 Properties of immiscible blends – design principles......................................................................29
3.2.2 Properties of miscible blends – design principles .........................................................................29
3.3 Block Copolymers and Fractal Structures........................................................................................31
3.4 Nanocomposites – a Route to New Materials..................................................................................34
3.5 Properties of Multicomponent Systems............................................................................................35
3.6 Dissolution and Mixing .....................................................................................................................37
3.7 Strategies for Dispersing Nanoparticles...........................................................................................39
3.8 Numerical Modelling.........................................................................................................................41
3.9 Advanced Functionality ....................................................................................................................44
3.10 Conclusions....................................................................................................................................47
3.11 References .....................................................................................................................................47
4 High-voltage applications of functional nanomaterials.......................................................................51
4.1 Introduction.......................................................................................................................................51
4.2 Cable ................................................................................................................................................52
4.3 Substation Equipment and Transformer..........................................................................................56
4.4 Rotating Machines............................................................................................................................63
4.5 References .......................................................................................................................................73
5 Emerging Regulatory Situations and Considerations of Nanomaterials for Dielectric
Applications...............................................................................................................................................77
5.1 Introduction.......................................................................................................................................77
5.2 Status & Emerging Regulatory Situations of Nanomaterials in selected Countries/Regions...........80
5.3 Perspectives to Reader / User of Nanomaterials in the Electrical Industry......................................88
5.4 Summary on Present and Emerging Regulatories of Nanomaterials ..............................................91
5.5 References .......................................................................................................................................93
6 Final remarks.....................................................................................................................................95

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EXECUTIVE SUMMARY
This brochure is the culmination of a formidable effort by a group of experts, integrating a wide body of
knowledge on nanodielectrics for electric power applications.
The last decade of research has demonstrated that it is possible to tailor dielectrics, thus opening the
possibility to implement functionality from a nanoscale modification or nanostructuration. The brochure
consolidates knowledge and illustrates the empirical basis for designing dielectric properties. Although some
guidelines are proposed for specific needs, further thoughtful reflections have resulted in a generalized
scheme, formalising the approach to designing dielectric properties from the nanoscale.
Nowadays, there are full-scale electrotechnical applications based on nanodielectrics. Instead of reviewing
existing applications involving nanodielectrics, this document will present a general analysis and review the
impact of nanomaterials, principally polymer nanocomposites, using specific examples. Nanotechnology
offers benefits for cables, transformer and substation equipment, and rotating machines. Advanced
nanocarbons in semiconductive shields are discussed. Multifunctional nanodielectrics for HVDC cable
application were considered in the context of space-charge control. Nanocomposite insulation benefits were
introduced for cast resin transformers. Moreover, current progress in nanodielectrics for fluid applications is
discussed. The use of functional nanomaterials in outdoor insulation has been mentioned in some details.
Some results from the European project ANASTASIA were presented and strategies for using
nanodielectrics in high-voltage rotating machines were mentioned.
Finally, introduction of new materials into products brings about a necessary reflection on the impacts of their
characteristics on the short, mid and long-term life and end-of-life of the specific applications. This exercise
has just started worldwide and it was judged worthy to present a comprehensive overview of the current
actions and trends involving regulations and reservations relative to the nanomaterials.
In conclusion the use of nanodielectrics has established itself as a potential and worthy solution in a variety
of practical cases involving electrotechnology.
1 Introduction
This working group (WG) follows the completion of the WG D1.24. In this latter working group, some
interesting progress was achieved. Common nanomaterial samples were prepared and shared among the
various participants, and a series of characterizations (micro and macro) were conducted. Many of these
results were reported [1.1 - 1.3]. After this focussed scientific action, it was judged appropriate to leave aside
further experimental work and consider future actions; the results of which are detailed in this report.
After 15 years of early research [1.4], there has been a continuous increase in scientific progress and
development in this field. Important sections of international scientific conferences are now devoted to
nanodielectrics and/or nanomaterials for electrotechnical applications. The research has demonstrated the
possibility of tailoring the macroscopic material properties at the nanoscale, and shown that the dielectric
properties could be affected and improved. These days we are exploring a second-generation of
nanodielectrics. With respect to these advanced materials, fabrication and processing take more prominence
in determining the properties exhibited by the prepared materials.
Furthermore, large research programs are being shaped and consolidated around the world on these topics.
For instance, Japanese researchers have been very busy exploring nanodielectrics and now they have
proposed a national initiative. China has developed very rapidly and has a national effort [1.5]. France is now
embarking on a national initiative dealing with “Supergrid” applications, that includes some implementation of
nanomaterials. Also, many developments were realized within the European 7th Framework Program (FP7),
e.g. the ANASTASIA project, and in the near future there will be the HORIZON 2020 Framework. In terms of
properties, end-users are expecting in the near future to have access to functional nanomaterials for
electrical power applications.
Nanomaterials are products of nanotechnologies. By definition, nanomaterials are materials that have at
least one dimension at the nanoscale (10
-9
m), e.g. nanoparticles, but also materials containing fillers at the

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nanoscale, e.g. nanocomposites. Physical and chemical properties of nanomaterials differ from those of
conventional materials.
Also, on October 18, 2011, the European Commission adopted the following definition [1.6] of a
nanomaterial: "A natural, incidental or manufactured material containing particles, in an unbound state or as
an aggregate or as an agglomerate and where, for 50 % or more of the particles in the number size
distribution, one or more external dimensions is in the size range 1 nm to 100 nm. In specific cases and
where warranted by concerns for the environment, health, safety or competitiveness the number size
distribution threshold of 50 % may be replaced by a threshold between 1 % and 50 %." This definition, still
debated, draws attention to the fact that industry and commerce could be further affected by the presence of
nanomaterials. Also, it adds constraints and requirements on the detailed characterization of the
nanomaterials.
In this study, nanodielectrics were almost focused on one category: polymer nanocomposites. Most of the
time, these consist of an inorganic phase having one nanometric dimension dispersed in a polymer matrix.
A functional material is a material able to perform a function. In that context, a multifunctional material would
be a material able to perform more than one function. In more general terms, the function can be a new or
enhanced material property. For instance, the electrical power industry seeks in the same dielectric the
simultaneous enhancement of breakdown strength together with higher thermal conductivity whilst
maintaining its mechanical performance.
This technical brochure contains 4 major chapters.
Chapter 1 is a brief introduction. Chapter 2 provides a snapshot of the current achievements.
Demonstrations of the abilities to tailor the functions of polymer nanocomposites are provided. Also, some
guidelines for developing functionality involving polymer nanocomposites are offered based on the last 15
years of acquired empirical knowledge.
Chapter 3 shows a more structured approach to preparing nanocomposites with the aim of exploring current
and new strategies for developing new functional nanomaterials and multifunctionality in polymer
nanocomposites.
Chapter 4 presents several examples of the potential use of polymer nanocomposites showing opportunities
to achieve multifunctionality. Several domains involving HV components and systems are illustrated, for
examples: cables, transformers, substation equipment and rotating machines.
Finally, Chapter 5 deals with the emerging regulatory framework of nanomaterials including environment,
health, safety, standardisation, registration and labelling. This chapter offers a worldwide view of those
evolving regulations and are classified by country.
[1.1] T. Tanaka, M. Fréchette, J. Kindersberger, S. Gubanski, A. S. Vaughan, S. Sutton, P. Morshuis, J-P. Mattmann, G. C. Montanari,
C. W. Reed, A. Krivda, J. Castellon, T. Shimizu, S. Pelissou, and M. Nagao, “Polymer nanocomposites,” CIGRE Technical
Brochure no. 451, pp.1–115, 2011.
[1.2] T. Tanaka, A. Bulinski, J. Castellon, M. Frechette, S. Gubanski, J. Kindersberger, G. C. Montanari, M. Nagao, P. Morshuis, Y.
Tanaka, S. Pelissou, A. Vaughan, Y. Ohki, C. W. Reed, S. Sutton, and S.-J. Han, “Dielectric properties of XLPE/SiO2
nanocomposites based on CIGRE WG D1.24 cooperative tests results,” IEEE Trans. Dielectr. Electr. Insul., vol. 18, no. 5, pp.
1484–1571, 2011.
[1.3] A. Krivda, T. Tanaka, M. Frechette, R. Gorur, P. Morshuis, S. Gubanski, K. Kindersberger, A. Vaugahn, S. Pelissou, Y. Tanaka,
L. E. Schmidt, G. Iyer, T. Andritsch, J. Seiler, and M. Anglhuber, “Characterization of epoxy microcomposite and nanocomposite
materials for power engineering applications,” IEEE Electr. Insul. Mag., vol. 28, no. 21, pp. 38–51, 2012.
[1.4] M.F. Fréchette, M. Trudeau, H.D. Alamdari and S. Boily, “Introductory remarks on nanodielectrics,” Proc. IEEE Conf. on Elect.
Insul. and Dielec. Pheno., Kitchener, Ontario, Canada, Oct. 14-17 (2001). An extended version was subsequently archived in
IEEE Trans. on Dielectr. and Electr. Insul., Vol. 11, No. 5, pp. 808-818, 2004.
[1.5] Shengtao Li, “Recent development on nanodelectric and its application prospect in China”, Workshop organized by the DEIS
Committee on nanodielectrics, Shenzen, China, October 20th, 2013.
[1.6] http://ec.europa.eu/environment/chemicals/nanotech/

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2 Some empirical guidelines to functionality
Since the introduction of the concept of nanodielectrics there has been an enthusiasm surrounding this new
field [2.1]. The majority of the last 15 years has been spent on examining polymer nanocomposite dielectric
behaviour. The research has tried to elucidate key parameters that play a role in the fabrication of
nanocomposites with the aim of tailoring them to desired properties.
Based on these research results and considering progress in other fields, this chapter proposes some
guidelines for the purpose of obtaining a dielectric property from a nanoscale manipulation. This is meant to
be a short account, not a review.
2.1 Basic Considerations
2.1.1 POLYMER NANOCOMPOSITE
By Polymer NanoComposite (PNC), it refers to a polymer matrix containing a nanometric sized filler. Most of
the observations discussed here concern a polymer matrix containing an inorganic phase. At times, remarks
may address PNC with nanometal, nanoceramic or semicon. A PNC does not exclude the presence of
several types of nanofillers to attain for instance a multifunctionality, these can be chemically bound (grafted
to polymer matrix) or not chemically bound to the polymer matrix.
The initial dimensions of the nanofiller and the final agglomerated state in the polymer are issues of
importance. The “appellation” is not strict yet and nanofillers having one dimension ≤ 100 nm are usually
considered [2.2, 2.3]. This limit is somewhat arbitrary and must be contemplated in the following context. It is
recognized [2.4] that the dielectric properties of polymer nanocomposites are dominantly determined by the
properties of the interfaces of the nanoparticle/polymer matrix. But, for the volume to be occupied by about
50 % of interface material, thus producing (viz. another assumption) a detectable macroscopic effect
resulting from a two-phase material, the assumption of an interface length of 10 nm restricts the diameter of
a rounded nanoparticle to 100 nm. Tanaka [2.5] argues that a loose layer could extend a few tens of nm –
but this region may be also considered the interphase and be associated with a morphology change. If taken
into account, it would push further the definition in relation to limits on dimensions.
Nanofiller agglomeration in the matrix is to be avoided. If there is agglomeration, it becomes difficult to
predict the outcome of this state as it may bring about the situation of having multiple defects or may
resemble a microcomposite-case. In general, the improved properties are linked to the physical presence of
the nanoparticles and by the interactions of the polymer with the particles and the state of dispersion.
2.1.2 SIZE RELATION WITH THE POLYMER
The literature [2.6-2.8] indicates that there is an intimate relationship developed between a polymer and the
nanofiller. Due to the high reactivity of the nanoparticle, the polymer dynamics are changed when
approaching the filler surface. Although the bulk structure does not change much, the polymer close to the
surface tends to change its morphology. This translates, for instance, into a change in the glass transition
temperature.
Adding filler particles to a polymer melt is expected to modify molecular arrangements and conformations of
the polymer chains on a global scale, and those with shorter segments of the polymer chains, on a more
local scale. The presence of the filler particles also modifies the conformational distribution of the polymer
chains with respect to the unfilled melt. There will be a favoured length that matches the spherical
nanoparticle and the length of the polymer chains, which falls in the range of 4 nm to 8 nm [2.9, 2.10].
Various fabrication techniques give access to this range, e.g. dissolution, sol-gel.
Attention was drawn by Fréchette [2.11] that to have a nanodielectric effect, a dielectric process at the
relevant scale must exist. It turns out that this relationship between scale and the nature of the phenomenon
is common. Quoted in Wikipedia [2.12] and briefly described, a study [2.13] published in Japanese, reports
many of these relations. The mechanical, electrical, thermal, optical, electrochemical, catalytic properties of
nanocomposites will be greatly different from those of the component materials. Size limits associated with

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these effects have been put forward: < 5 nm for catalytic activity, < 20 nm for making a hard magnetic
material soft, < 50 nm for refractive index changes, and < 100 nm for getting superparamagnetism,
mechanical strengthening or restricting matrix dislocation movement.
2.1.3 COMPATIBILITY
In case of organic/inorganic nanocomposites, as mostly considered for nanodielectrics, the strength or level
of interaction between the organic and inorganic phases are important. Simple mechanical mixing causes
only weak dispersion and low interactions. To form strong interactions (hydrogen bridges, van der Waals
forces, chemical covalent or ionic bonding) other processes are required. Sol-gel technique, in-situ
polymerization, melt compounding by twin-screw extrusion, solution blending, high-shear mixing or ball
milling are suitable preparation methods for organic/inorganic nanocomposites [2.14-2.17]. Andritsch et al.
reviewed the influence of manufacturing on the dielectric performance of nanocomposites [2.18]. Besides the
preparation method, the chemical nature of the organic polymer matrix and the inorganic nanofiller are of
crucial importance for a good compatibility.
For nanodielectris, the diversity of the polymer matrices ranges from thermoplastics, e.g. polyethylene (PE),
thermosets, e.g. epoxy resin, to rubbers which may be unpolar. This is summarized in Table 2.1. To achieve
good compatibility between nanofillers and the matrix polymer, the polarities of the matrix and the filler have
to be aligned. For an improvement in the dispersion of particles in the matrix, the simple addition of a
surfactant can be sufficient to reduce cohesion between the particles. This surfactant builds an interphase
between the particles and the matrix, but there is no chemical link. This can of itself lead to an improvement
in properties, but with limits, since such surfactants may also lead to a weakening of the polymer matrix.
Poly-
ethylene
Poly-
styrene
Polyester Epoxy resin
uncured/cured
Polarity Low low medium high/medium
Molecular
weight
High high High low/infinite
Crystallinity high high medium low
Monomer
formula
uncured
cured
Table 2.1: Important traits of the polymer matrices

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An effective way to couple nanoparticles and polymers is to apply a chemical surface modification of the
nanoparticles, which is described in detail in section 2.2.2. Surface modification also improves the dispersion
of the particles in the matrix as surfactants do by reducing the cohesion between the particles and increasing
the adhesion (hydrogen bondings, van der Vaals forces) but additionally enables a covalent bonding
between the matrix and the nanoparticles. Considering the chemical structure of the polymers, it is
necessary that for a PE-matrix a different modification to the nanoparticles is required than for epoxy resins,
also discussed in section 2.2.2. Appropriate surface modification of the nanoparticle leads to improved
properties of the nanocomposite [2.19].
2.1.4 PERMITTIVITY: THE POLYMER OR THE ADDITIVE
Besides worrying about the polarity match between the nanofiller and the polymer matrix, one could ask if
the nature of the dielectric permittivity of the additive may be used to tailor the permittivity of the composite.
For an additive having a dielectric constant ~10, e.g. alumina (Al2O3), it can be demonstrated that the
resulting dielectric constant of a microcomposite containing as much as 50 wt% can remain as low as the
Polyethylene (PE). A demonstration [2.20] illustrated in Fig. 2.1 shows clearly that the preparation technique
affects the dielectric characteristics, yet when adequately prepared (here degassed and ball-milled), the
resulting permittivity of the composite appears almost flat as a function of the excitation frequency and close
to that of PE.
For nanocomposites [2.21], irrespective of the polar nature of the matrix (PE vs epoxy), it is found that it is
rather the role of the absorbed water at the interface that appears to be a dominant factor in affecting the
dielectric response [2.22, 2.23]. This behavior is thought to arise from the interplay between interfacial water
mobility and bonding rather than the host polymer. Obviously, this will be more present when “oxides” are
involved.
Incidentally, the real permittivity of a polymer nanocomposite can be largely enhanced by starting with a high
dielectric-constant polymer and adding some special nanoceramics (large weight, e.g. titanates) or
nanoconductors (above percolation) [2.24]. Applications include gate dielectrics, high charge-storage
capacitor and electroactive materials. Obviously low dielectric losses are also required.
Figure 2.1: Real permittivity of PE-Al2O3 films obtained
from different mixing procedures (after [2.20])

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2.1.5 LARGER OR SMALLER WT% CONTENT
In connection with their large surface-to-volume ratio, the rationale for the use of nanofillers is that much less
active material would be required. However, there are caveats and variants on this argument. From the
several ranges discussed in literature, in general, at least for dielectric applications, a wt% range below 5
would be reasonable.
Beyond this content, re-agglomeration and/or interactions between the nanoparticles or defects tend to bring
a detrimental effect. Agglomerates were well documented for the case of an epoxy nanocomposite
containing 20 wt% of nanosilica [2.25]. On a macroscopic level, the gain obtained from using POSS
(Polyhedral Oligomeric Silsesquioxanes) added to epoxy [2.26] in its surface resistance to erosion was found
to start to decline when 1 wt% is exceeded. This is exemplified in the next figure.
Figure 2.2: Evaluation of resistance to electrical discharge by comparison
of eroded sample volume, after [2.26]
Other non-electrical properties are found to vary also based on wt% reaching a maximum at various low
wt%. For instance, the thermal stability of polymers with incorporated metal nanoparticles [2.27] often
exhibits very peculiar characteristics. After the expected improvement in thermal stability through the
incorporation of 2 wt% of Cu nanopaticles into low-density polyethylene, further increasing the amount of Cu
nanoparticles leads to a deterioration in the thermal stability.
In summary, there is no ideal wt%, it depends on the polymer matrix, the nature, size and shape of the filler
and the type of the property to be affected.
2.2 Fillers
What type of nanofillers should be used to attain a certain dielectric property? There is no simple answer to
this question, there are a multitude of documents providing ample information on nanofillers, consult some of
the following, e.g. [2.3, 2.28 – 2.31]. Some comments are offered below.
2.2.1 BASIC FEATURES
Initial state
The nanofillers would come as an agglomerated powder. In several instances, it is better to acquire or
prepare them in a solvent or as a colloidal solution. A micrometric phase made of agglomerated
nanostructures may show different macroscopic properties. In 2001, Fréchette et al. [2.32] gave the example
of the “nanovaristance” effect, associated with large grains stemming from sintering nano-additives.

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Dimension
At least one dimension should be in the range of 1 nm to 100 nm. It does not imply that the dimension is
distributed between 1 nm and 100 nm: some effects require that the size distribution be narrow. Many times
the nanometric feature is associated with rather extravagant dimensions. For instance, nanoclay and
graphene oxide may exhibit lengths of hundreds of microns. By the way, a filler with a dimension attaining
120 nm would be a nanofiller.
Shape
Nanofillers come as particles, platelets, tactoids, rods, wires, flakes, tubes, whiskers, etc. Of course, the
surface-to-volume ratio will change with the shape, with more interface there is more reactivity. Flat or
platelet shapes are found to be effective for increasing the surface resistance to erosion while also affecting
fire-retardant performance and improving gas barrier effect. Orientation effects may be possible with
elongated shape-factor particles. For strictly insulating purposes, shapes bringing about local electric-field re-
enforcement should be avoided.
Nature
The important thing is to evaluate purity. Some natural impurities (e.g. in clay) or fabrication residues may
strongly affect the outcome. Available nanofillers are numerous and fall into many categories: inorganic,
organic, metallic, semi-conductor, etc.
Dispersion
The dispersion of the nanofiller within the whole of the matrix is a key factor in achieving success.
Functionalization
In most cases, the best results are obtained when the surface of the fillers is functionalized. A basic
treatment is to use silane. The compatibilization treatment will be a function of the chemistry of both filler and
polymer matrix.
Insulating
The metallic oxides, e.g. TiO2, SiO2, fumed silica, and the nanoclays are very popular. Non-oxides like BN
and AlN are gaining use, which are also known for their high thermal performance.
Semiconductor
Silicium Carbide already used as grading-field materials have their nano-sized counter parts. Carbon blacks
contain already nano-sized particles. Nano-ZnO [2.33] and Graphene Oxide (GO) [2.34] are certainly
candidates for future applications. The aim for future applications is to reduce greatly the wt% needed to
attain percolation-threshold field conductivity. Instead of 60 wt% to 30 wt%, we are speaking of a fraction of
that wt%. In some circumstances, Carbon NanoTubes (CNT) may be considered semi-conducting.
Conductor
There are various nano-metals available, see for instance [2.35]. These additives can be used in affecting
the electrical conductivity of a polymer. Furthermore, they often come with an insulating coating, natural or
deposited as in the case of nanorod of gold surrounded by silica.
POSS (Polyhedral Oligomeric Silsesquioxanes)
There is a renaissance in this field where some early active developments were taking place in the sixties
[2.36]. Polyhedral oligomeric silsesquioxanes (POSS) may be thought of as a nanometric form of silica. One
of their most common forms (Si8O12R8) consists of a cubic cage of eight silicon corner atoms and twelve
oxygen edge atoms, where each of the eight silicon atoms carries one of an extremely wide spectrum of
functional groups (R), allowing hundreds of possible compounds. In their properties, polyhedral oligomeric
silsesquioxanes occupy a middle ground between silica and polysiloxanes. POSS is not a nanoparticle, it is

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considered a hybrid organic/inorganic with nanometric size. It can be pictured as a silica box surrounded by
functionalization. It does not agglomerate but may form large polycrystalline structures. Reactive POSS will
form covalent bonds with polymers. POSS used in conjunction with Polypropylene (PP) was reported by
Takala et al. [2.37]. Composites formed by PP and POSS behave as dielectrics. Their research reported on
some substantial improvements relative to dielectric behaviour. Other research [2.38] involving PE confirmed
a mild improvement in dielectric performance, e.g. the breakdown strength. However, change in morphology
with low wt% content may be of some consequence in some future applications.
Extra-ordinary fillers
There are nanofillers that exhibit marvelous intrinsic properties. Graphene is such an example. Graphene is
a semi-metal with a high electron mobility at room temperature and thermal conductivity in the range of
5 × 103
W m−1
K−1
. However, its properties are not directly transferable in the process of manufacturing a
polymer nanocomposite (see for instance [2.39]). Generally speaking, it is still difficult to transfer the amazing
performance of these nanofillers to the polymer composites.
Nanofiller into polymer
The fabrication techniques will affect greatly the outcome relative to macroscopic properties. There are the
chemical routes (e.g. dissolution [2.40], grafting [2.41]), mechanical ways (e.g. ball-milling [2.42], mixer
[2.43], extruder) and several others to produce polymer nanocomposites. The approach selected has been
shown to be an important factor shaping the final properties [2.43].
2.2.2 CHEMISTRY AND FUNCTIONALIZATION
Nanofillers for electrical insulation applications are basically inorganic materials which are non-conductive
and have a certain thermal conductivity: Nanoclays, Metal oxides (Alumina oxide, Titanium oxide, etc.), Non-
Metal oxides (Silica oxide particles, Phytosilicates) and Nitrides (Boron nitride, Alumina nitride).
Besides their morphology and shape (layered, non-layered, particle shape, L/D-ratio, tubes, 3-dimensional
structures, etc.), they have a quite common surface chemistry. All oxides possess hydroxyl-groups on their
surface which are open for chemical modification as shown in Figure 2.3. Even though nitrides are non polar,
they also possess a sufficient amount of hydroxyl-functionality on the surface, so that they can be modified in
the same manner as oxides without any pre-treatment to introduce oxide-functionality [2.44]. Furthermore,
for a good dispersion of layered materials, an exfoliation of phyllosilicates and clays is essential to break the
adhesion between single layers.
Figure 2.3: Scheme of surface chemistry of oxides and nitrides
Metal or Non-metal
oxide particles

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As mentioned in section 2.1.3, one major issue with the preparation of nanocomposites is the dispersion of
the nanoparticles and the creation of a strong interface between the particles and the matrix. This is of vital
importance for nanodielectrics to unleash their full potential [2.45, 2.46].
Surface modification of nanoparticles using silanes has turned out to be a most effective way. The basic
structure of the silanes is R-(CH2)n-SiX3, wherein R represents the functional group which is able to form a
bond or another chemical link to the matrix and -(CH2)n- represents a spacer. X can be either chlorine or an
alkoxy-group. In the silanization reaction (Equation 1), X is hydrolyzed in the first step to hydroxyl
Equation 1: First step of the silylation reaction: Hydrolyzation
groups which are condensing in the second step with the hydroxyl groups of the nanoparticle surface and/or
with other hydroxyl groups (Equation 2).
Equation 2: Second step of the silylation reaction
This principle scheme is common for the silanisation of all particles. The surface is changed from –OH to –R
and the ability of the surface modified nanoparticle to interact and form bonding to the matrix is now
determined by R. Table 2.2 represents a choice of silanisation agents according to the structure R-(CH2)n-
SiX3. If Polyethylene is the matrix, C2H5- or CH2=CH- residues can be selected to achieve compatibility or
even covalent bonding. If epoxy is the matrix, R can be an amine, anhydride or glycidyl group.

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Table no. 2.2: Coupling agents for surface modification of nanoparticles, after [2.47]
2.3 Guiding Considerations
2.3.1 PREAMBLE TO DESIGNING FUNCTIONALITY
Based on the final application, a polymer may be chosen from one of the three groups: thermoplastics,
thermosets and/or rubbers. Polymer properties can be adjusted or modified by tuning: the formulations
(fillers, additives, plasticisers, stabilizers, pigments) and/or the processing (solvent methods, in situ
polymerization, melt blending, sol-gel method).
A polymer nanocomposite can be defined as a polymer matrix into which nanoparticles (nanofillers) have
been introduced to form a blend. For the nanocomposites, particles sizes, shape and composition, nanofiller
concentration, can be varied. Among the most common nanoparticles are nanoclays, carbon nanoparticles
(MWNT, SWNT, graphene, etc.), silica and metal oxides. Since the nanoparticles have much higher surface
area than microparticles, special attention is given to the interactions or interface zones. One of the most
efficient ways to engineer the interface regions is to chemically adjust the particle surface to modify its
bonding with the matrix. A modified surface will help the dispersion of the nanoparticles in the matrix. The
processing technique and parameters used to produce the nanocomposite can change the dispersion of the
fillers in the polymer matrix and so the final properties. The morphology, glass transition temperature Tg, and
free volume of polymers can provide important indirect information about the nanocomposite interface [2.31].
The addition of fillers can enhance a certain property of the matrix, such as Young’s modulus, thermal
stability and/or introduces new properties such as magnetism, thermal and electrical conductivity, etc. These
property improvements concern volume properties (modulus, strength), surface properties (hardness,
abrasion resistance), dimensional stability, thermal stability, as well as optical and electrical properties. The
composites also exhibit reduced permeability and chemical stability (UV resistance).
When dimensions of inorganic fillers are reduced to nanoscale, much less filler is needed to achieve superior
performance. The incorporation of only a few percent of nanosized particles can make property changes. An
example of the enhancement in mechanical properties shown in Fig. 2.4.

Page 13
Figure 2.4: Enhancement in mechanical properties of polymer nanocomposites:
(a) PDMS/silica nanocomposite (in situ polymerization); (b) Neat NR and composites with
4 phr of carbon black, clay fibrous, clay layered and MWNT [2.48]
Presence of rigid particles results in enhancement of the elastic modulus of a polymer composite compared
to that of the neat polymer. The reinforcement can be caused in principle by one of the three main
mechanisms: (i) the stress transfer from the matrix to the non-isometric particles [2.49], (ii) the partial
substitution of soft matrix with a stiffer filler [2.49] and (iii) the segmental immobilization caused by the
interaction of polymer chains with filler surface [2.50]. The stress transfer mechanism depends on the
inclusion aspect ratio, its orientation to the applied load and the strength of adhesion; it is considered as a
size independent contribution for sufficiently flexible matrix chains. The second mechanism is independent of
particle size as well. On the other hand, the segmental immobilization mechanism contributes to the overall
composite reinforcement with the extent primarily affected by the size of the inclusions, becoming important
for submicron particles with large surface to volume ratio [2.51].
Depending on the strength of the interaction between polymer and particle, the interaction zone can have a
higher or lower mobility than the bulk material, and result in an increase or decrease in Tg. It has also been
suggested that free volume in such interaction zones is altered by the introduction of nanofillers. Since, these
interaction zones are likely to overlap at relatively low volume fractions in nanocomposites, a small amount of
nanofiller has been found to impact the electrical behaviour [2.30, 2.52].
Although the term nanodielectric can mean any dielectric system on a nanoscale, the term nanodielectric is
frequently used to refer to polymer/nanoparticle mixture or nanocomposite. Lewis [2.52] emphasized that the
interaction zone around the particles is a “quasi-conductive” region which partially overlaps in the
nanocomposites. These overlapped interface regions thus may allow charge dissipation, which, in turn, could
be expected to improve the dielectric breakdown strength and voltage endurance characteristics [2.31, 2.53].
2.3.2 MODIFICATION OF THE DIELECTRIC RESPONSE
The dielectric response of PNC, nano-sized metallic oxide and layered silicate filled polymers, has been
investigated to a considerable extent and reported in a number of publications [2.21, 2.53, 2.54-2.57]. In
addition to dipolar relaxation mechanisms proper to each phase and induced from a possible compatiblizer
polar group, such as maleic anhydride, interfacial polarization is usually detected, but with a dielectric
strength and a relaxation time largely depending on the quality of the dispersion and/or the presence of
absorbed moisture. For hydrophilic fillers, absorbed water, particularly in the interfacial regions between
fillers and the polymer matrix, is almost inevitable and will invariably affect the composite dielectric response,
often leading to an interfacial relaxation peak at a much higher frequency than what would be expected
based on values of the filler bulk conductivity. This effect is expected to be even more pronounced for PNC
as the volume fraction of interfaces is much higher and it has been clearly shown in the case of silica-based
PNC [2.56,2.58]. As a result of water absorption at the particle-matrix interface, a displacement of the

Page 14
relaxation frequency to higher frequencies has been observed with increasing thickness of the interlayer
coating. This gives a surprising result: the frequency of maximal dielectric loss shifting to lower values with
increasing temperatures as seen in Fig. 2.5(a) for the case of an untreated nanosilica-filled polyethylene.
When silica particles are treated with a coupling agent, the absorption of water is considerably reduced.
However, the dielectric loss in the dry state might be higher in the case of the treated particles due to the
dipolar nature of the coupling agent itself, as shown in Figure 2.5(b). A more detailed investigation of this
particular effect can be found in [2.56].
(a) (b)
Figure 2.5: (a) imaginary permittivity as a function of frequency for PE filled with 2 wt%
nanosilica [7]; (b) imaginary permittivity of treated (FN) and untreated (UFN) nano-silica
polyethylene composites before (UAN) and after (AN) annealing at 60°C for 7 days
A number of polymers can be reinforced by layered silicate fillers and another intriguing dielectric response is
the one obtained from layered silicate reinforced polymers. Master batches of these PNC are now
commercially available and dilution can be easily conducted by melt compounding either with an extruder or
a mixer. Various degrees of exfoliation or intercalation can be reached depending on the compounding
process, the chemical formulation of the organically-modified clay platelets, and the use of coupling agents,
or compatibilizers.
Figure 2.6 illustrates the variation of the imaginary permittivity as a function of frequency for the two different
PNC systems, with (Fig. 2.6b) and without (Fig. 2.6a) compatibilizer. A main relaxation peak moving towards
high frequencies can be clearly observed in both cases. In the case of the PNC, with little or no exfoliation,
see Figure 2.6(a), two relaxation processes are observed at 60o
C. This double peak dielectric response has
been previously reported [2.57, 2.59, 2.60], with the high frequency relaxation peak being assigned to a
dipolar relaxation process [2.57, 2.60], while the low frequency relaxation peak has been explained by an
interfacial relaxation process.
1.0E‐05
1.0E‐04
1.0E‐03
1.0E‐02
1.0E‐01
1.0E‐02 1.0E‐01 1.0E+00 1.0E+01 1.0E+02 1.0E+03 1.0E+04 1.0E+05 1.0E+06
Imaginary Permittivity
Frequency (Hz)
25oC ‐ start
30 oC
40oC
50oC
60oC
70oC
80oC ‐ First run
80oC ‐ Second run
25oC ‐ end

Page 15
(a)
(b)
Figure 2.6: Real and imaginary parts of the permittivity for temperatures from 0°C to 60°C;
(a) 3 wt% LLDPE/clay nanocomposite; (b) 3 wt% LLDPE/clay nanocomposite
with 10% coupling agent [2.55]
Figure 2.7 compares the dielectric loss for the exfoliated sample, as compared to the less ex-foliated sample.
One can see that the predicted lowering of the interfacial relaxation rate with the change of the filler aspect
ratio [2.59] is also observed experimentally.
2
2.5
3
3.5
4
r
'
10
-2
10
-1
10
0
10
1
10
2
10
3
10
4
10
5
10
6
10
-4
10
-3
10
-2
10
-1
10
0
Frequency (Hz)
r
'' LLDPE nClay3%,T=0°C
LLDPE nClay3%,T=20°C
2
2.5
3
3.5
4
r'
10
-2
10
-1
10
0
10
1
10
2
10
3
10
4
10
5
10
6
10
-3
10
-2
10
-1
10
0
Frequency (Hz)
r''
LLDPE nClay 3+10%,T=0°C

Page 16
Figure 2.7: Imaginary permittivity at 60°C for PE/clay nanocomposites with
and without compatibilizers
For a possible future application in power cables, only the dielectric response at power frequencies 50/60 Hz
really matters. Table 2.3 below gives the dissipation factor at temperatures of 20°C, 40°C and 60°C, for
different PE/clay nanocomposites and at power frequency. Despite the tremendous increase of the dielectric
loss due to the addition of nanoclays (two orders of magnitude), the dielectric losses still remain low enough
to cause no concern for power cable applications. Indeed, for a typical distribution cable insulation of 12 mm
inner radius and 19 mm outer radius, operating at a phase-to-ground voltage of 14.4 kV, and dissipation
factor of 10-2
, the power loss would be less than 1 W/m, i.e. at least one order of magnitude less than the
Joule heating loss of the conductor at full load.
20°C 40°C 60°C
LLDPE 3.59 x 10-4
3.75 x 10-4
9.74 x 10-4
LLDPE/0/1 8.10 x 10-3
1.33 x 10-2
8.81 x 10-3
LLDPE/0/3 2.14 x 10-2
4.35 x 10-2
3.59 x 10-2
LLDPE/10/3 5.12 x 10-3
3.59 x 10-2
4.49 x 10-2
LLDPE/0/5 3.88 x 10-2
8.53 x 10-2
8.88 x 10-2
Table no. 2.3: Dissipation factor at power frequency (60 Hz) for LLDPE/nanoclay
composites [2.59]
2.3.3 MODIFICATION OF THE DIELECTRIC BREAKDOWN
In 2005, Dongling Ma et al. [2.19] studied LDPE/TiO2. They were the first to demonstrate the great
importance of the properties of the interface on affecting the breakdown of a polymer nanocomposite.
Obviously, the interface must be functionalized to attain a larger breakdown value as compared to that of the
neat polymer. Fig. 2.8 shows the results from several experiments reported by Roy et al. [2.61]. It is shown
clearly that if the interface is treated appropriately that the breakdown of the nanocomposite may exceed that
of the XLPE.

Page 17
Figure 2.8: Breakdown behavior of XLPE/silica nanocomposites, with the specificity of the
interface as a parameter
Furthermore, the work of Roy et al. was even more dramatic for voltage endurance tests, where 2.5 orders of
magnitude improvement were found for nano-SiO2 in XLPE using the coupling agent tri-ethoxy-vinylsilane
versus XLPE alone. The dramatic role of specific coupling agents was confirmed, since the improvement
was not found for nano-SiO2 in XLPE without the use of a coupling agent or by the use of a different coupling
agent. The bonding of the three ethoxy groups to hydroxyl groups on the nanoparticles resulted in an
increase of 40% in crystallinity and a dramatic drop in surface free volume, were attributed to these effects.
At the recent conference on solid dielectrics (ICSD 2013), several avenues for the self-healing of polymer
nanodielectrics when subjected to ac and dc electric stresses, during testing or operational service are
identified and discussed [2.48]. These include a) the application of (controlled/current limited) so-called
“clearing” dc breakdown stresses in the factory following the manufacture of metalized industrial, HED, and
monolithic film capacitors); and b) the use of electron scavenging additives, stress grading additives,
expandable monomers (to negate void formation), and of sacrifial nanocapsules, for use with solid extruded
cables, insulation for use in rotating machines, and with solid insulation for aerospace, aircraft, and medical
systems use.
2.3.4 MODIFICATION OF SURFACE RESISTANCE TO DISCHARGES
Electrical insulation is subjected to surface partial discharges (PDs) in high voltage apparatus in many
occasions. Therefore PD resistance is our long-time main concern, and is always expected to be increased.
Nanotechnology may help. It was found that PD resistance is enormously enhanced by the addition of small
amount of nanofiller to polymer matrices. PD resistance is usually evaluated by IEC (b) electrode system
and/or a rod-to-plane electrode system. Epoxy is mainly studied [2.62 – 2.70] but other materials such as
polyethylene and polypropylene are also partly investigated [2.71– 2.72]. Possible mechanisms are
discussed in the reference [2.66].
First of all, it should be pointed out that good and homogeneous dispersion of nanofiller in a polymer of
interest is a requisite. As a data example, Figure 2.9 indicates that PD resistance is affected by the kind of
fillers. Possible mechanisms, still hypothetical though, are proposed on the basis of experimental results that
have been obtained to-date, as shown in Figure 2.10. These are useful as a guideline to modify insulating
polymers to obtain highly PD resistant nanocomposites.

Page 18
Figure 2.9: Comparison of PD resistance of epoxy nanocomposites loaded with different
kinds of nanofillers
1. There is a positive effect of nanofillers on PD resistance. Both particle-like fillers such as silica and
titania and belt-like fillers such as layered silicate have a similar effect.
2. Tight interfaces and well-ordered morphology will increase PD resistance. Silane couplings and
grown spherulites are some of the examples for that.
3. Two and three dimensional segmentation of organic polymers by inorganic fillers seems to work well
against the material attack by partial discharges.
4. Stacking nanoparticles will suppress progress of PD erosion: Extrinsic effect.
Organic resin parts are weak against PDs. Inorganic filler parts are strong against PDs. Nano- segmentation
divides resin at the nanoscale. Strong interfaces mean the increase in diameter of filler and the decrease of
resin regions.
On the basis of the proposed mechanisms, PD resistant enamel wires and generator winding insulations are
developed [2.73-2.74].

Page 19
Figure 2.10: Various factors creating more PD resistant properties
of polymers with the aid of nanofillers
2.3.5 MODIFICATION OF ELECTRICAL CONDUCTIVITY
It is possible to modify the electrical conductivity of polymer using nanostructuration. There are several ways
and, especially, several ranges of conductivities of interest. However, in most applications, the polymer must
remain in an insulating state. This leaves a span of about two orders of magnitude for the conductivity to
vary.
A good example is provided by Jonscher in his book [2.75]. The measurements involve PA6/MWCNT, PA6
being Polyamide 6 and MWCNT, multi-walled carbon nanotube, respectively. These polymer nano-
composites were produced by compression molding. AC conductivity was measured using Broadband
Dielectric Spectroscopy. A transition from insulating to conductive state can be observed in Fig. 2.11. This
transition is seemingly related to the conducting filler and its connectivity (percolation). The transition is
identified by the appearance of dc conductivity, i.e. ac conductivity becomes independent of frequency. The
plateau values give the dc conductivity. The transition is seen to occur between 2.5 wt% and 5.0 wt% CNTs.
With increasing carbon nanotube content, dc conductivity was found to increase and the plateau to extend to
higher frequencies. At each composition, the frequency dependence of ac conductivity can be seen in
Fig. 2.11.

Page 20
Figure 2.11: AC conductivity as a function of excitation frequency in PA6/MWCNT
One of the advantages of using some nano-additives is to produce a sharp transition from insulating to
conducting state using only a fraction of filler wt%. One example that involves the evaluation of dc
conductivity in terms of percolation with nanocomposite of Polycarbonate (PC) containing MWCNT, follows
[2.76]. Percolation theory for a nanocomposite with conducting inclusions (like CNT) in an insulating matrix
(PC) is driven by the equation: σdc(p) ~ (p- pc)t
where p is the volume fraction of conducting inclusions (see
inlet, Fig. 2.12). Pc is the percolation threshold and t a fractional power law exponent depending on the
dimensionality. Fig. 2.12 shows experimental results. The line is the fit of the percolation equation to the
data. The inlet shows the same results yet represented differently. The percolation threshold is found to be
much lower than the prediction based on a 3-dimensional statistical distribution of conducting spheres (about
16 vol%, [2.77]). This is most probably due to the elongated shape of the carbon nanotubes.
Figure 2.12: Evaluation of dc conductivity in terms of percolation;
p is the filler concentration in wt%, after [2.76]
Nano-metals can also be used to tailor the conductivity. For instance, nano-silver was found to make vary
the electrical conductivity over 4 orders of magnitude [2.78, 2.79].

Page 21
2.3.6 MODIFICATION OF THERMAL PROPERTIES
Thermal properties in electrical insulation are mainly concerned with (1) thermal expansion, (2) thermal
conductivity and (3) thermal endurance. These performances can be modified by adding suitable nano and
micro fillers. Firstly, thermal expansion coefficient of insulation should be matched to that of metal when
insulation is in contact with metal conductors. Switchgear apparatus is a typical example. Microfillers such as
silica and alumina are required for that purpose. In order to obtain good electrical performance under this
condition, nanofillers are further added into the microcomposite resulting in a nano-microcomposite.
Successful results have been obtained relative to the development of compact switchgear insulation [2.80].
Figure 2.13: Thermal conductivity of epoxy/hexagonal-BN composites
Secondly, thermal conductivity is increased in general by heavily loading of microfillers with high thermal
conductivity because its enhancement requires percolation of fillers inside polymer matrices. A target value is
something like 10 W/m/K. This value can be obtained by using BN for example [2.81]. Dielectric breakdown
strength is generally decreased due to inclusion of microfillers, and nanofiller such as silica is recommended
to be added [2.82]. Nanofillers can also modify the thermal conductivity itself if it is intended to modify it in a
certain lower range up to 1 W/m/K [2.83].
Inorganic fillers such as alumina, aluminum nitride (AlN), boron nitride (BN), silicon carbide and even
diamond are usually utilized to increase thermal conductivity of polymers [2.84]. A value larger than
10 W/m/K can be obtained by using alumina and AlN if they are highly loaded even up to more than
90 vol.%. Void formation makes it difficult to obtain a target value in reality. An example of composites with
less filler content is shown in Figure 2.13 [2.85]. Various types of epoxy/BN composites were prepared to
optimize material conditions for high thermal conductivity and high breakdown strength. The best condition is
obtained for composites like NChM-5, i.e. epoxy/conglomerated h-BN/nanosilica nano-micro-composite. In
general the loading of macro- fillers will reduce breakdown strength of a composite, but the further addition of
nanofillers will increase once lowered value again [2.86]. Useful recipe is suggested as shown below:
1. Loading inorganic microfillers with high thermal conductivity results in high thermal conductivity composites.
Further addition of nanofillers will produce better electrical insulation.
2. Suitable orientation of fillers is required if a filler of interest has anisotropy in thermal conductivity.
Conglomerated BN filler is one of the examples. Orientation with the electric field will help.
3. Reduction of void content is one of the most important factors to prepare composites from the standpoint of the
intended increase in both thermal conductivity and breakdown strength. In that sense, co-mixing with different-
sized fillers and hybrid mixing with different shape fillers are useful techniques to reduce void formation.
4. Surface treatment and nanofiller addition will contribute to additional increase of thermal conductivity and
breakdown strength.

Page 22
Thirdly, thermal endurance is expected to increase in some cases. Glass transition temperature Tg is an
important parameter to evaluate this performance. It was found that Tg increases in an interaction zone, i.e.
a region between a filler particle and a matrix, in epoxy and other materials [2.87]. Some of the examples are
shown in Figure 2.14.
Another example is shown in Figure 2.15 [2.88] for silicone rubber with silica to improve tracking resistance.
Thermal ablasion by laser is used for this evaluation. The following message is useful as a guideline to
produce heat resistant nanocomposites:
1. Scanning electron microscopy analysis of silicone rubber composites prepared from various nanofillers with
Triton has shown that the surfactant greatly improves the dispersion of nano-sized particles, yielding
nanocomposites that are more homogeneous, with improved resistance to heat ablation.
2. The surfactant appears to be beneficial to the dispersion, as long as the surface covered by the surfactant is
lower than the BET surface area of the fillers. Fumed silica was shown to impart greater heat ablation resistance
than either natural silica or alumina.
3. There is no big difference in the erosion resistance of natural silica- or alumina-filled compositions. The ablation
observed on nano silica-filled specimens suggests that the silica accumulated at the surface forms a heat-
resistant barrier preventing further erosion of the underlying silicone rubber. This phenomenon was not
observed in the alumina filled specimens.
4. Composites consisting of an admixture of micron-sized and nano fumed silica display significantly improved
resistance to heat ablation than compositions with only one or the other filler, particularly when Triton is used to
disperse the particles. This suggests a promising area of study for outdoor insulation applications.
Figure 2.14: Interlayer spacing (film thicknesses) in model nanocomposites that yield the
same Tg (glass transition temperature) deviation as 0.4 vol% silica-PMMA and silicaP2VP
nanocomposites. Tg deviation of P2VP model nanocomposites (open squares) and
PMMA model nanocomposites (open circles). Right and left: Transmission electron
microscopy images of 0.4 vol% silica-P2VP (right) and 0.4 vol% silica-PMMA (left)
nanocomposites (scale bars = 100 nm). The error bars ( 1 K) represent the inherent
error due to the fitting of the data required to obtain Tg

Page 23
Figure 2.15: Relationship of the additive amount of TRITON and the amount of erosion in
silicone rubbers with 2.5 wt% fumed silica (colloid silica)
2.4 Closing Remarks
The last decade of research involving PNC has opened the gate to designing dielectrics at the nanoscale. In
this chapter, accumulated knowledge and information on PNC has been structured so as to construct a
guideline for modifying relevant macroscopic properties and/or attaining improved performance. A selection
of properties have been addressed. Others like space charge accumulation, treeing resistance, electrical
endurance, etc. were omitted for the sake of brevity. However, they can also be affected, controlled or
enhanced. The overall picture presented here is much in phase with the state-of-the-art of the research field.
Finally it is concluded that the way is open to the use of multiple additives to control several properties at the
same time and for the dielectrics to be truly multifunctional.
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3 Strategies towards novel nanostructured functional
dielectrics
3.1 Introduction
As technology advances, the requirements placed on materials of all types increase. As far as polymers are
concerned, the worldwide rate of consumption is rising at about 5 % per annum as a result of increased
usage in established sectors and through the use of polymer-based systems as the materials of choice in
new areas. In automotive applications, for example, the adoption of advanced materials, such as main chain
liquid crystal polymers (LCP), leads to weight reduction benefits in concert with the ability of the material to
survive harsh environments for the life of the vehicle. As in other classes of polymers, the properties of main
chain LCPs are strongly influenced by the chemistry of the basic units that form the polymer chain. To
illustrate this, contrast the electrical characteristics of two polymers based upon carbon and hydrogen,
namely, polyethylene (PE) and polyacetylene (PAc) [3.1]. The ethylene monomer can be represented
CH2=CH2, which leads to an idealised polymer backbone structure that is a fully saturated sequence of CH2
groups joined by carbon/carbon single bonds; a very good insulator. In polyacetylene, the polymer backbone
is made up of a sequence of CH groups that can be seen to be connected together, at one extreme, by an
alternating sequence of carbon/carbon single and double bonds. However, in reality, the resulting π electron
cloud is delocalised to a degree along the molecule, to give a polymer that can exhibit electrical
conductivities comparable with that of copper. So, one way to produce polymers with new functionalities that
can be used in new applications is to synthesise new monomers and polymerise them; the problem with this
approach is primarily economic, in that it requires major investments in manufacturing plant and the resulting
specific cost of the material will remain high until the price is pushed down by increasing consumption. The
alternative strategy to the development of systems with improved properties or increased functionality is to
combine existing materials in new ways.
1. Additives. Antioxidants, for example, are added to polymers as a direct means of enhancing thermal
stability during high temperature processing or during long service lives, while clarifiers such as
dibenzylidene sorbitol (DBS) affect the optical properties of the system indirectly. That is, DBS forms
a highly disperse gel phase at concentrations above ~1000 ppm, the gel induces massive nucleation
in the polymer and it is this morphological change in the base polymer that reduces scattering and
gives a final material with high optical clarity. The area of small molecular additives will not be
considered here; for more information see, for example, the review by Pfaendner [3.2].
2. Blends. Blending together different polymers constitutes a powerful and well established route to the
formation of “new” materials with improved properties. In general, such systems can be categorized
as miscible or immiscible, depending upon the tendency of the constituent components to mix or
segregate under pertinent conditions. Both of these systems will be considered here, together with
the related topic of block copolymers. While block copolymers are not strictly blends, they certainly
contain chemically distinct elements, each of which contributes to the final behavior of the system
and, as such, they have certain similarities with immiscible blends, albeit that phase separation is
limited by covalent bonding between the blocks.
3. Composites. This strategy is well exemplified in the commercial aerospace sector by the
replacement of metals by composites in the Airbus A380 and Boeing 787 Dreamliner. Elsewhere,
much of the current interest in nanocomposites is often cited as originating as a result of research
conducted at Toyota [3.3], which showed how the performance of thermoplastics such as
polyamides could be improved by introducing a very small quantity of a suitable nanoclay, such that
the resulting system could be employed in new, more demanding applications. That is, through the
addition of the nanofiller, the functionality of the material is increased.

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3.2 Polymer Blends – Functionality through Mixing
In the examples that follow in this section, modified functionality is derived by combining different polymers.
Consequently, it is convenient to deal with the issues in terms of miscibility.
3.2.1 Properties of immiscible blends – design principles
In immiscible blends, the two polymeric components form distinct phases that can be readily identified
structurally and where each confers some desirable characteristic on the final system. The classic example
of this concerns rubber toughened amorphous polymers, such as high impact polystyrene (HIPS).
Polystyrene (PS) is a clear, amorphous thermoplastic, with a glass transition temperature, Tg, around 115°C
and, as such, at room temperature, it exhibits typical glassy mechanical properties of brittleness and a low
elongation at break of ~ 4 %. Consequently, in isolation, it is unsuitable for many applications, where a tough
material is required. To improve the functionality of the polymer in this respect, a second rubbery phase can
be included, but since conventional rubbers interact only weakly with PS, there is no thermodynamic drive for
mixing at the molecular level (see section 3.6). Consequently, a dispersed, second phase is formed.
Typically, to produce such a system, the rubbery component of the blend is dissolved in the styrene
monomer before polymerization and, to achieve the maximum improvement in toughness, both the size of
the rubbery particles and their interaction with the matrix need to be controlled. 10 % to 20 % of the rubber
phase is generally added in the form of latex particles 0.1 µm to 1.0 m in size and, to minimize interfacial
failure, they are first coated with an appropriate moiety that results in the formation of covalent bonds
between the two blend phases. The result is a material that retains the attractive processing characteristics
of PS but in which the elongation at break is increased by up to an order of magnitude and the impact
strength by a factor of ~3, albeit at the expense of a small decrease in ultimate tensile strength and softening
temperature. Another polymer where an equivalent strategy has been successfully employed is
polycarbonate (PC), where blends with materials such as rubbery acrylonitrile/butadiene/styrene triblock
copolymers can give materials with excellent impact performance combined with good optical clarity and
improved high temperature stability compared with HIPS.
3.2.2 Properties of miscible blends – design principles
The influence of polymer morphology on the electrical properties of polymers has been studied for many
decades. In 1980, Kolesov [3.4] published a key paper on the effect of spherulite size on the breakdown
strength of polyethylene and polypropylene. This work reported that, as the spherulite size increased, so the
breakdown strength decreased, with the implication that optimum breakdown performance would be
obtained from a material containing no large scale morphological features. That is, a material rather like
crosslinked polyethylene (XLPE). In the early 1990s, Vaughan and co-workers at the University of Reading
began to re-examine this idea, by growing spherulites in blends of high crystallinity high density PE (HDPE)
and low crystallinity low density PE (LDPE), where the LDPE formed the majority component (typically
~ 80 %). These workers showed that, if crystallization were conducted isothermally at a temperature where
the HDPE could crystallise but the LDPE could not, it was possible to generate materials with significantly
improved breakdown strengths [3.5] and resistance to electrical treeing [3.6]. The explanation for this
enhancement in properties centres on the sequential crystallization process that occurs in such a miscible
binary blend, whereby, a skeleton of relatively thick crystals composed, primarily, of HDPE initially forms
within a matrix of liquid LDPE. Only on cooling, does the LDPE crystallise and this occurs uniformly within
the pre-established HDPE framework. The consequences of this two-stage crystallization process is that the
mechanisms that gave rise to the weak inter-spherulitic boundaries observed by Kolesov [3.4] and elsewhere
are, effectively, engineered out of the system. The final materials were also found to be flexible at low
temperatures and to exhibit good high temperature mechanical integrity, presenting the possibility of being
used at higher temperatures than conventional XLPE insulation. Subsequent work has demonstrated that
cables can be manufactured from such HDPE/LDPE thermoplastic blends.
While it is convenient to discuss blending in terms of miscible and immiscible, this is something of an
oversimplification and, in reality, it is necessary to consider the processes by which phase separation may
occur. First, phase separation may occur in the melt such that, as in the case of the amorphous systems
described above, a two-phase morphology is generated (so-called liquid-liquid phase separation).
Alternatively, phase separation may occur during crystallization, as a consequence of one component of the

Page 30
blend crystallising preferentially (so-called liquid-solid phase separation). This will inevitably lead to changes
in the residual melt composition which, themselves, can initiate further liquid-liquid phase separation. The
resultant crystalline morphology can then be complex and will reflect both the crystallization itself and the
effect of phase separation in the residual melt.
An example of this in the context of material functionality and dielectrics relates to a natural extension of the
work on polyethylene blends described above. An ideal material for high voltage alternating current (HVAC)
cable applications would exhibit the following characteristics:
 High breakdown strength
 Low electrical conductivity
 Low dielectric loss
 High thermal conductivity
 High thermal stability
 Good flexibility across a wide temperature range
While polyethylene meets many of these requirements, the maximum operating temperature of any PE-
based cable will be limited in such a way as to act as a significant restriction on cable ratings; a way around
this is, therefore, to employ an insulation system based upon polypropylene. However, isotactic
polypropylene (iPP) is a stiff material that is brittle at low temperatures, while copolymers of propylene and -
olefins are generally characterized by relatively low melting temperatures and, consequently, offer few
advantages over PE. A way to address such shortcomings is through blending, whereupon the iPP can be
thought to act in a similar way to the HDPE in the system described above and an appropriate propylene-
based copolymer is used instead of the LDPE. While this general concept has long been appreciated, as can
be demonstrated by reference to the copious patent literature in the area, actually designing a system that
will self-assemble under manufacturing conditions to give the required electrical and thermo-mechanical
functionality is a significant challenge [3.7]. Put simply, the problem centers on how to incorporate a sufficient
quantity of a flexible copolymer to generate a mechanically acceptable material without compromising
breakdown strength. In the case of polypropylene blends, miscibility between the homopolymer and many
copolymers is limited. Consequently, phase separation of the least crystallisable molecular fractions within
the copolymer tends to occur such that, unless crystallization occurs extremely quickly, a discrete and
defective second phase tends to form, as shown in Fig. 3.1. While this is positively beneficial in terms of the
impact performance of the immiscible blends described above, the electrical implications are unacceptable.
Consequently, designing a material which exhibits the overall functionality that is required is problematical.
Figure 3.1: Scanning electron micrograph showing the phase separated structure of a blend
of isotactic polypropylene and a propylene-based copolymer containing 12 % ethylene
co-monomer (scale bar 10 m)

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3.3 Block Copolymers and Fractal Structures
3.3.1 BLOCK COPOLYMERS
A key feature in the examples given above concerns optimisation of the phase structure and interfacial
interactions between the various constituents in a blend. A natural extension of the immiscible blend concept
is therefore to include the required functionalities within a single molecular architecture, as in the case of
block copolymers. Continuing with the theme of designed mechanical materials, the family of
acrylonitrile/butadiene/styrene (ABS) copolymers provides a good example of this in that each of the three
molecular blocks within the structure is included to bring a particular macroscopic characteristic to the final
material: the styrene block provides a rigid phase; the butadiene block provides a rubbery phase; the
acrylonitrile, being relatively polar, tends to increase interchain interactions. However, in the case of
materials such as ABS, it is not just the block properties but also the block length and the resulting material
morphology that determine the final material properties. Block copolymers exhibit a rich variety of phase
structures and, in the case of an AB diblock system, increasing the fraction of A from 0 % to 50 % results in a
change in local structure from one based upon a cubic array of spherical inclusions of A within a continuous
matrix of B, through a hexagonal array of cylinders of A within a continuous B phase through a bi-continuous
gyroid structure to a lamellar texture. The theory of structural evolution in such materials has attracted a
great deal of attention and has been relatively well understood for many years, at least in the case of simple
systems (e.g. see [3.8]). In terms of properties, appropriate tailoring of the block structure can lead directly to
desirable characteristics and, in the case of ABS, for example, polymerizing styrene and acrylonitrile in the
presence of polybutadiene to give a system containing 15 % to 35 % acrylonitrile, 5 % to 30 % butadiene
and 40 % to 60 % styrene results in a material with excellent shock absorbing characteristics.
3.3.2 PERCOLATION AND FRACTAL STRUCTURE
The examples described above have focussed on mechanical properties. However, it is also possible to
modify electrical characteristics by combining electrically conducting and insulating components and, as in
the above examples, the precise phase structure that results is critically important in determining ultimate
bulk properties.
The classic means of modifying the bulk conductivity of an insulating polymer is through the addition of a
conducting carbon filler and, as a result, carbon black filled polymers have long been used in electrical
applications such as cable screens. The general topic of assemblies of discrete conducting fillers within an
insulating matrix has therefore attracted considerable experimental and theoretical attention; here, we
provide an overview of electrical conduction – the thermal analogue is considered in rather more detail
below. For polymer filled with carbonaceous particles, dynamic percolation is described in the literature as a
self-assembly process that occurs in the quiescent melt [3.9-3.12]. It is evidenced by a tremendous
conductivity increase similar in nature to that of the statistic percolation transition [3.13]. In a recent study, for
example, Huang et al. [3.14], considered a wide range of factors relating to processing of polyethylene filled
with carbon black. These included the effect of processing conditions on molecular parameters, thermal
characteristics, morphology and electrical conductivity. This study revealed that, in all cases, the variation of
conductivity with filler loading level followed a classical percolation behaviour in which the conductivity of the
composite increased from values in the range 10-18
S cm-1
to 10-16
S cm-1
for carbon black contents of 0 wt%
to 20 wt%, to more than 10-4
S cm-1
at 30 wt% carbon black and above. Since no compositions between
20 wt% and 30 wt% filler were considered in this work, the precise form of the percolation onset is unclear,
but an increase in electrical conductivity of twelve orders of magnitude over a composition range of just
10 wt% of filler is, nevertheless, dramatic. Comparable experimental results are also cited in the work of
Yang and Liang [3.15], who adopted an effective medium approach to the problem of electrical conduction
through carbon black filled polymers, in which they assumed that the key conducting element was not a
single carbon black particle, but rather, aggregates of these primary particles. The resulting theory predicted
an extremely sharp percolation threshold, in which conductivity increased by some 12 orders of magnitude
over a composition range of about 1 vol%; comparison with experiment suggests that, while this may be
rather more extreme than is seen in practice, the essential features are nevertheless reproduced.

Page 32
Although carbon black has long been the most widely used form of carbon in such applications, the
availability of other more exotic nanostructured forms (fullerenes, carbon nanotubes, graphene, etc.) has led
to the recent study of composites based upon these materials and, while many of the basic principles still
apply, significant differences result from changes in nanofiller aspect ratio.
To illustrate this, consider the effect of introducing carbon nanotubes (CNT) into a thermoplastic matrix. In
the case of systems based upon PC and multi-walled CNT (MWNT) produced by melt compounding from a
master batch [3.16], the final dispersion was found to be critical in determining properties and, therefore,
factors such as mixing time and screw speed exerted a significant influence on the observed percolation
threshold. Nevertheless, in contrast with the particulate systems described above, for these materials, only
1 wt.% to 2 wt% of the MWNT was required to switch the material from insulating (conductivity ~10-16
S cm-1
)
to conducting (conductivity > 10-4
S cm-1
). Al-Saleh et al. [3.17] reported on another study of materials
including PC and CNT, but where polystyrene was also included to form an immiscible blend as the matrix.
In this system, the spatial location of the CNTs was found to be dependent upon the CNT loading level; at
low CNT concentrations (~0.05 wt%), the CNTs were found to be located preferentially in the PS phase
whereas, at higher levels (~5 wt%), they were uniformly distributed throughout both polymer phases. When
processed appropriately, the tendency for the CNTs to be concentrated in the PS led to the formation of a
conducting network within the system at a CNT content of just 0.05 wt%. An alternative approach that has
been employed in an attempt to force the formation of contiguous conducting structures involves mixing
CNTs with ultra-high molar mass polyethylene (UHMWPE) [3.18]. In this case, the authors effectively coated
the granules of UHMWPE with CNTs before hot compacting these to form what they termed “a segregated
and double percolated structure”. That is, the CNTs form a locally percolated structure on the surfaces of the
PE granule surfaces which, themselves, provide macroscopic percolation paths, such that macroscopic
electrical conduction was achieved at just 0.049 vol% of CNT.
Another class of material that has been shown capable of exhibiting unconventional forms of electrical
behaviour as a consequence of the formation of network structures are blends of insulating and conducting
polymers. For example, Reghu et al. [3.19] observed that the percolation limit in a blend of polyaniline (PANi)
and poly(methyl methacrylate) (PMMA) occurred at just 0.3 % of PANi. The reason for this can be seen in
Fig. 3.2, which shows a transmission electron micrograph obtained from a similar blend. From this, it is
evident that the blend morphology is based upon a bi-continuous phase structure, in which the PANi (dark
structural features in the figure) adopts a fractal structure that is nanometric in scale [3.20]. This could,
therefore, be considered as a self-assembling material in which the mechanical properties are largely
determined by the majority matrix phase and the electrical conductivity by the percolating fractal structure of
the PANi. In materials such as this, the overall functionality reflects the various components of the system
and, by appropriate control of both, together with the interactions that occur between them, it is possible to
modify the response of the overall system. Figure 3.3 shows how it is possible to use this strategy to
generate robust functional materials with conductivities that span the complete range from insulators through
to moderate conductors [3.21].
Figure 3.2: Bright field transmission electron micrograph showing a blend of polyaniline
camphor sulphonic acid (3.8 %) in poly(methyl methacrylate) (scale bar 1 m)

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