FABRICATION AND CHARACTERIZATION OF MULTIFUNCTIONAL
POLYETHERIMIDE/CARBON NANOFILLER COMPOSITES
A dissertation submitted in partial fulfillment of
the requirements for the degree of
DOCTOR OF PHILOSOPHY
WASHINGTON STATE UNIVERSITY
School of Mechanical and Materials Engineering
© Copyright by BIN LI, 2012
All Rights Reserved
© Copyright by BIN LI, 2012
All Rights Reserved
To the Faculty of Washington State University:
The members of the Committee appointed to examine the dissertation of Bin Li find it satisfactory
and recommend that it be accepted.
Weihong (Katie) Zhong, Ph.D. (Chair)
Mohamed Osman, Ph.D.
Louis Scudiero, Ph.D.
Jin Liu, Ph.D.
As the date of my dissertation defense is getting closer, I am proud to say going to WSU for my
PhD study is the best decision I have ever made in past few years. WSU has an open-minded,
innovative, supportive and enthusiastic research environment with great opportunities. I always
believe that individual success is really limited, and we can harvest more in a superior research and
working environment. In addition, I appreciate that college of engineering and architecture, and
school of mechanical and materials engineering for providing such a great academic environment
for me and every other graduate student and helping us achieve success.
I truly enjoy my graduate student life with my advisor, Dr. W. H. Katie Zhong. It is my pleasure
to work in her research group in past four years. She is always supportive of my enthusiasm in
research and encourages me to adventure different opportunity in academic communities. She is
not only a good advisor who is always helping me figure out my problems, but also a mentor of my
life and always gives me valuable suggestions.
I need to thank all my current and past graduate committee members: Dr. Mohamed Osman, Dr.
Louis Scudiero, Dr. Jin Liu as well as Dr. Hussein Zbib. Thanks for your time, patience as well as
generous contributions to my research project and PhD study.
I appreciate all my intelligent and kind-hearted graduate group mates: Weston Wood, Brooks
Lively, Tracy Ji, Tian Liu, Zack Tang, Brady Beacon and Allen Eyler. It is memorable experiences
to work with all of you. Your intelligence and generosity always enlighten me in both of my
academic and personal lives. I also thank all the undergraduate students who have worked with me:
Loren Baker, Anthony Perugini and Erik Olson. They gave me a lot of young and interesting
memories, and I also learned so much from them.
At last, I also acknowledge all funding agencies and companies who support my PhD research
project: NSF (GOALI Grant 0758251and NIRT Grant 0506531), the Boeing Company, Sabic
Innovative Plastic Co., Applied Science, Inc. as well as Cytec Industries Inc. (Cytec Engineered
FABRICATION AND CHARACTERIZATION OF MULTIFUNCTIONAL
POLYETHERIMIDE/CARBON NANOFILLER COMPOSITES
by Bin Li, Ph.D.
Washington State University
Chair: W.H. Katie Zhong
Polyetherimide (PEI) is a high-performance thermoplastic, showing great potential in many
important applications including airplanes and electronics, due to its excellent mechanical and
thermal properties. At the same time, with the rapid development of nanotechnology, adding
various nanomaterials to polymeric materials to fabricate polymer nanocomposites has become a
popular and efficient way to develop advanced materials. Graphitic carbon nanofillers (GCNs),
benefiting from their unique SP2 hybridized carbon structure, exhibit superior mechanical, thermal
and electronic properties, and prove to be an all-purpose nanomaterials for next-generation polymer
This project focuses on development of new generation PEI/GCN nanocomposites with
substantially improved mechanical properties, static dissipation and acoustic damping; as well as
functionalities to satisfy a variety of applications including airplanes, ground transportation and
electronics. In order to achieve these goals, efforts have been made to control dispersion as well as
surface modification of GCNs in PEI matrix. Dispersion of GCNs is a critical factor for properties
of polymer nanocomposites, impacting both fundamental mechanical properties and physical
properties. At the same time, proper surface modification can improve dispersion of GCNs,
strengthen interfacial bonding between PEI and GCNs, and impart new functionalties to the
resultant nanocomposites. The results showed success in comprehensively improving properties and
functionalities of PEI/GCN nanocomposites. Effective control of GCN dispersion and efficient
surface modification has substantially improved mechanical, thermal, tribological, damping, and
static dissipation properties of PEI/GCN nanocomposites. The structure-property relationships of
PEI nanocomposites have also been discussed in detail for further materials design and optimization
of material properties. Based on these relationships, a novel non-destructive evaluation method has
been developed to quantitatively examine the dispersion of GCNs in polymer nanocomposites. This
research may substantially expand potential applications of PEI and its nanocomposites.
TABLE OF CONTENTS
ACKNOWLEGEMENTS ................................................................................................................................. iii
ABSTRACT ....................................................................................................................................................... v
CHAPTER 1 Introduction .................................................................................................................................. 1
1.1 High performance polymers and polyetherimide ............................................................................... 1
1.2 Nanomaterials and graphitic carbon nanofillers (GCN) ..................................................................... 3
1.3 Functionalties of polymer / GCN composites .................................................................................... 7
1.3.1 Mechanical properties ..................................................................................................................... 8
1.3.2 Tribological properties ................................................................................................................... 9
1.3.3 Electrical properties ........................................................................................................................ 9
1.3.4 Dielectric properties ..................................................................................................................... 13
1.3.5 Damping properties ...................................................................................................................... 15
1.3.6 Other properties ............................................................................................................................ 16
1.4 Dispersion and distribution of GCNs in polymers ........................................................................... 19
1.4.1 Dispersion and distribution control during composite Processing ............................................... 20
1.4.2 Improvement of dispersion via surface modification ................................................................... 21
1.5 Problem statement ............................................................................................................................ 23
1.6 Research objectives and significances .............................................................................................. 25
References ..................................................................................................................................................... 28
CHAPTER 2 Single Negative Metamaterials in Unstructured Polymer Nanocomposites toward Selectable
and Controllable Negative Permittivity ............................................................................................................ 34
Abstract ......................................................................................................................................................... 34
2.1 Introduction ...................................................................................................................................... 35
2.2 Experiments ...................................................................................................................................... 37
2.2.1 Raw materials ............................................................................................................................... 37
2.2.2 Preparation of polymer nanocomposites ...................................................................................... 37
2.2.3 Morphology analysis .................................................................................................................... 38
2.2.4 Dielectric and electrical tests ........................................................................................................ 38
2.3 Results and discussion ...................................................................................................................... 38
2.3.1 Physical origins and concentration dependence of negative permittivity ..................................... 39
2.3.2 Effect of morphologies of CNFs ................................................................................................... 45
2.3.3 Effect of chemical structures of polymers .................................................................................... 46
2.4 Conclusions ...................................................................................................................................... 48
References ..................................................................................................................................................... 50
CHAPTER 3 The Strong Influence of Carbon Nanofiber Network Variablity on the Pronounced AC
conductivity of the Polyetherimide Composite Films ...................................................................................... 52
Abstract ......................................................................................................................................................... 52
3.1 Introduction ...................................................................................................................................... 52
3.2 Experiments ...................................................................................................................................... 54
3.2.1 Materials ....................................................................................................................................... 54
3.2.2 Control of configuration of CNF network .................................................................................... 55
3.2.3 Microstructure analysis ................................................................................................................. 56
3.2.4 Electrical properties analysis ........................................................................................................ 57
3.3 Results and discussion ...................................................................................................................... 57
3.3.1 Electrical properties under DC and AC electrical fields ............................................................... 57
3.3.2 Microstructures and correlations with electrical properties .......................................................... 60
3.3.3 Mechanism analysis ...................................................................................................................... 62
3.4 Conclusions ...................................................................................................................................... 66
References ..................................................................................................................................................... 67
CHAPTER 4 Effectual Dispersion of Carbon Nanofibers Polyetherimide Composites and Their Mechanical
and Tribological Properties .............................................................................................................................. 69
Abstract ......................................................................................................................................................... 69
4.1 Introduction ...................................................................................................................................... 70
4.2 Experiments ...................................................................................................................................... 73
4.2.1 Preparation of polyetherimide/carbon nanofiber composites ....................................................... 73
4.2.2 Three point flexural testing ........................................................................................................... 74
4.2.3 DMA measurement ....................................................................................................................... 74
4.2.4 Sliding wearing test ...................................................................................................................... 74
4.2.5 Microstructure analysis ................................................................................................................. 75
4.3 Results and discussion ...................................................................................................................... 76
4.3.1 Dispersion of carbon nanofibers in PEI matrix ............................................................................ 76
4.3.2 Flexural property analysis............................................................................................................. 79
4.3.3 Dynamic mechanical properties ................................................................................................... 84
4.3.4 Sliding wear properties ................................................................................................................. 87
4.4 Conclusions ...................................................................................................................................... 90
Reference ...................................................................................................................................................... 92
CHAPTER 5 Effective Static Dissipation of Bi-layer Thermoplastic Nanocomposites at Low Nanofiber
Loadings ........................................................................................................................................................... 94
Abstract ......................................................................................................................................................... 94
5.1 Introduction ...................................................................................................................................... 94
5.2 Experiments ...................................................................................................................................... 98
5.2.1 Raw materials ............................................................................................................................... 98
5.2.2 Preparation of polymer nanocomposites ...................................................................................... 98
5.2.3 Morphology and properties analysis ............................................................................................. 99
5.3 Results and discussion .................................................................................................................... 100
5.3.1 Electrical and dielectric properties of mono-layer composites ................................................... 100
5.3.2 ESD protection performance of mono-layer composites ............................................................ 104
5.3.3 Bi-layer composites with improved ESD protection performance at low loadings .................... 105
5.3.4 Comparison of ESD between bi-layer and mono-layer composites ........................................... 110
5.4 Conclusions .................................................................................................................................... 111
References ................................................................................................................................................... 113
CHAPTER 6 Simultaneous Enhancements in Damping and Static Dissipation Capability of Polyetherimide
Composites with Organosilane Surface Modified Graphene Nanoplatelets .................................................. 115
Abstract ....................................................................................................................................................... 115
6.1 Introduction .................................................................................................................................... 116
6.2 Experiments .................................................................................................................................... 118
6.2.1 Materials ..................................................................................................................................... 118
6.2.2 Silanization of graphene nanoplatelet ........................................................................................ 119
6.2.3 Preparation of PEI/GNP nanocomposites ................................................................................... 121
6.2.4 Fourier transform infrared spectroscopy .................................................................................... 121
6.2.5 Dispersion and interface analysis ............................................................................................... 122
6.2.6 Dynamic mechanical analysis..................................................................................................... 122
6.2.7 Electrical and dielectric properties ............................................................................................. 123
6.2.8 Thermal analysis ......................................................................................................................... 123
6.3 Results and discussion .................................................................................................................... 123
6.3.1 FTIR ........................................................................................................................................... 123
6.3.2 Stability of PEI/GNP solution .................................................................................................... 124
6.3.3 Morphology of PEI/GNP nanocomposites ................................................................................. 127
6.3.4 Dynamic mechanical /damping properties ................................................................................. 128
6.3.5 Static dissipation property .......................................................................................................... 136
6.3.6 Thermal stability ......................................................................................................................... 139
6.4 Conclusions .................................................................................................................................... 140
References ................................................................................................................................................... 142
CHAPTER 7 Effect of Non-Covalent Surface Modification via Poly(3,4-ethylenedioxythiophene)-
poly(styrenesulfonate) on Electrical Properties of Porous Polyetherimide/Carbon Nanotube Nanocomposites
Abstract ....................................................................................................................................................... 144
7.1 Introduction .................................................................................................................................... 144
7.2 Experiments .................................................................................................................................... 149
7.2.1 Materials ..................................................................................................................................... 149
7.2.2 Surface modification of CNT ..................................................................................................... 149
7.2.3 Preparation of PEI/PEDOT:PSS/CNTnanocomposites .............................................................. 150
7.2.4 Examination of surface modification and dispersion analysis .................................................... 150
7.2.5 Electrical properties .................................................................................................................... 151
7.3 Results and discussion .................................................................................................................... 151
7.3.1 Effect of non-covalent dispersion on CNT dispersion ................................................................ 151
7.3.2 Microstructures of porous PEI/PEDOT:PSS/CNT hybrid composites ....................................... 155
7.3.3 Anisotropic volume electrical properties of PEI hybrid nanocomposites .................................. 156
7.3.4 Surface electrical properties ....................................................................................................... 159
7.4 Conclusions .................................................................................................................................... 161
References ................................................................................................................................................... 163
CHAPTER 8 Conclusions and Future Plans .................................................................................................. 165
APPENDIX .................................................................................................................................................... 168
Project 1 Novel Hydration Induced Flexible Sulfonated Poly(etherketoneketone) Foam with Super
Dielectric Charateristics .............................................................................................................................. 168
Project 2 High Modulus Aliphatic Polyimide From 1,3- Diaminopropane and Ethylenediaminetetraacetic
Dianhydride: Water Soluble to Self - Patterning ........................................................................................ 194
1.1 High performance polymers and polyetherimide
High performance polymers are a group of aromatic polymers with extraordinary mechanical
properties (high strength and modulus), superior thermal / thermal-oxidative stability and chemical
resistance. The important HPPs include polyimide (PI), Poly(aryletherketone)s (PAEKs), Poly (p-
phneylene sulfide) (PPS), polysulfone (PSU), Polycarbonate (PC) etc., which are playing an
increasingly important role in our daily life, compared with commodity polymers [1-5]. The
significant applications of these high performance polymers include electronics, packaging,
structural materials for aerospace and automotive, capacitor, battery, fuel cell and medical implants,
etc [6-20], thus, they have gained growing interests in aerospace, national defense and many
industrial applications, such as electronics, medical and transportation.
It is well recognized that the polymer chain structures consisting of covalently bonded phenyl
groups and strong inter /intra-molecular interactions (hydrogen bond, dipole-dipole force, van der
waals force, etc.) are the fundamental structural factors accounting for the high performances of
these polymers, especially, high mechanical and chemical stabilities at high temperature. However,
while these two structural factors lead to the superiority of high performance polymers over other
commodity polymers, their negative side–effects are also obvious: it is really difficult to mold them
into desired morphologies and shapes via melt processing technique, due to extremely high melt
viscosities as a result of rigid polymer structure and poor chain mobility.
Polyetherimide (PEI) is a derivative of polyimide (PI) materials with flexible ether group
introduced onto rigid and highly polar imide chain structures, in order to improve the melt
processiblity while maintain the outstanding properties of PIs. Its chemical structure is given in
Scheme 1.1. Some basic properties of PEI and PI are summarized in Table 1.1. When there is no
significant difference between of PEI and PI in basic mechanical performances, the processing of
PEI is much easier, according to much lower Vicat softening temperature as well as processing
temperature. In addition, PEI also has a low dielectric constant of 3.15, compared with PI materials.
Scheme 1.1 Chemical Structure of Polyetherimide (PEI)
Table 1.1 Comparisons between polyetherimide (PEI) and polyimide (PI)
1120 36500 5 218 350-400
1050 34800 4 260 380-410
Note: these data are provided by Sabic Innovative Plastic Inc.
The better processibility and its superior properties make PEI a good replacement of PI in many
applications, including aerospace, ground transportation, structural materials as well as electronics.
However, the practical applications of PEI still face some critical problems. Firstly, the further
improvement of basic properties is still needed, such as strength and modulus, in particular, its
brittleness and poor wear resistance; secondly, contrary to the urgent need of polymer based organic
energy materials and electronics, especially those requiring high temperature stability, the
functionalities of PEI have been rarely explored.
1.2 Nanomaterials and graphitic carbon nanofillers (GCN)
Given the extremely small size and consequent large number of surface free electrons in
nanoparticles (Fig. 1.1) , the optical, electronic, and magnetic properties, as well as the
chemical reactivity, of nanomaterials are distinct from those of microparticles (MPs) and bulk
materials , as a result of quantum-size effect at the nanoscale. The range of 1–100nm is the most
accepted interval of measurement for nanomaterials [22, 23].
Figure 1.1 Figure 1.1 Distribution of micro- and nano-scale fillers of the same 0.1 vol.% in a
reference volume of 1 mm3 (A: Al2O3 particle; B: carbon fiber; C: graphite Nanoplatelet; D:
Carbon nanotube) 
The burgeoning nanotechnology gives birth to fruitful innovative nanostructures and
nanomaterials  showing distinguished and unique functionalities which directly depend on the
chemical structures, morphologies and size of nanomaterials, for example, the distinct light
scattering properties of nano-gold and nano-silver particles with different sizes and morphologies
as given in Fig.1.2. This feature makes them very useful in bio-diagnostics.
Figure 1.2 Sizes, shapes, and compositions of metal nanoparticles can be systematically varied to
produce materials with distinct light-scattering properties 
Graphitic carbon nanofillers (GCN), including graphene, graphite nanoplatelet (GNP), carbon
nanofibers (CNF), carbon nanotubes (CNT) and fullerene, have proved to be a group of useful
functional nanomaterials, consisting of graphenes with a honeycomb network of sp2 hybridized
carbons (Fig.1.3). In other words, all kinds of graphitic carbon nanofillers can be obtained by
manipulating graphene monolayers. Graphene does not only have extraordinary mechanical
properties including high Young’s modulus (~1100GPa), fracture strength (~125GPa); but also have
excellent physical properties including extremely high thermal conductivity (~5,000Wm−1
mobility of charge carriers (200,000cm2V
−1), and extraordinary transport [26-33]
Figure 1.3 One atom thick graphene monolayer and the conversion of graphene into other graphitic
nanostructures (From left to right: fullerene, CNT and GNP)
properties as a result of its unique band structure as shown in Fig.1.4. As early as 1947, Wallace
theoretically studied the electronic properties of graphene and his results showed that graphene has
a unique band structure in which the conduction and valence bands just touch each other, forming
an exactly zero-band gap semiconductor , which have been proved by Dr. A.K.Geim’s group
who firstly reported an experimentally isolated single-layered graphene in 2004. The energy
dispersion relation of the two bands is therefore linear in wave vector, k, and they cross at the Dirac
point of the two-dimensional Brillouin zone (Fig.2). Close to the Dirac point, the graphene
dispersion relation takes the form as given in equation 1.1:
where c*is the velocity of charge carriers indicating that graphene is a zero-gap semiconductor with
symmetric bands (Fig.2C) . This property, in combination with the hexagonal crystal lattice
symmetry, makes the electrons behave as if they were massless fermions, governed by Dirac’s
equation, showing new form of chiral quantum Hall effect  as well as fractional quantum Hall
Figure 1.4 (A) Brillouin zone of graphene with two inequivalent lattice points, K and K’. (B) Linear
dispersion relation of graphene, forming Dirac cones above and below the Dirac point . (C)
Approximation of low-energy band structure of graphene 
The outstanding properties of graphene are also imparted to graphitic nano carbons, resulting in
their outstanding mechanical, thermal, electrical and optical properties, etc [38-50] (Table 1.2), the
applications of which include electronic devices, photonics, capacitor, battery, electrical discharging
protection, electro-magnetic shielding, acoustic damping, solid lubricant, and perhaps biomedical
applications such as drug delivery and controlled release, etc. Thus, they are overwhelmingly
considered as ideal candidates for reinforcing and functional components in new generation
(B) (A) (C)
Table 1.2 Some properties of GNP, CNT and CNF
Diameter (nm) Aspect ratio
(W /m K)
GNP ~5 (thickness) N/A ~105 ~3000 ~1000 -
CNT ~1 (SWCNT)
~10(MWCNT) >1000 ~10
3 ~ 4000 ~1200 ~30
CNF ~100 ~200 ~10-3
~2000 ~600 ~ 7
1.3 Functionalties of polymer / GCN composites
Although some fascinating characteristics of polymers, such as low mass density, flexibility,
elasticity, conformability, good processibility and unique physical/chemical properties, make them
promising in new generation smart materials, structural materials as well as energy materials, etc.,
the practical applications are still limited without proper modification, due to low mechanical and
thermal stabilities and inadequate functionalities in most polymers. By compounding polymers and
nanomaterials, not only can the basic properties be dramatically improved, but also many significant
functions could be successfully realized [38-45].
A great number of works have suggested the superiority of GCNs in high performance and
functional polymer nanocomposites. There are no other nanomateirals like GCNs that could
comprehensively improve basic properties of polymers, while achieving new functions, due to
extraordinary mechanical and physical properties of GCNs in electronics, energy as well as
biomedicine, etc., as demonstrated in Section 1.2. The success of GCN reinforcement in polymeric
materials includes dramatically increased mechanical, tribological and thermal properties, as well as
tailorable electrical and dielectric properties and so on.
1.3.1 Mechanical properties
GCNs possess extra high modulus and strength, as shown in Table 2, compared with polymeric
materials, whose modulus and strength are usually only several GPa and far below 1GPa,
respectively. With very small loadings of GCNs, the load bearing capability of polymers can be
significantly improved, leading to better modulus and strength of polymer matrix [51-53].
Meanwhile, via proper interface modification, and dispersion control of nanofillers, the carbon
nanofillers could raise the fracture toughness of the polymer matrix by activating/enhancing relative
motion of polymer chains [54-56]. Besides the dramatically improved mechanical performances,
another attractiveness of GCNs is only very small amount of GCNs is needed, compared with
micro-size carbon fillers, which could effectively main the low mass density, flexibility and
transparency of polymers. Thus, it is getting more and more popular. In the studies of epoxy/CNF
nanocomposites , only 0.3wt% loading of surface modified CNFs could effectively increase the
flexural strength, flexural modulus, and fracture toughness by 26%, 21% and 36%, respectively, as
a result of uniform dispersion of CNFs as well as strong interfacial bonding between surface
modified CNFs and epoxy matrix. The flexural properties of epoxy could also be comprehensively
improved by the reinforcement of ball milled GNPs with the loading as low as 0.5wt% . The
addition of 0.2wt% CNFs with proper surface modification also significantly increased the tensile
fracture toughness of ultra high molecular weight polyethylene by 16 times . However, the
existence of nano particles in a polymer matrix also creates a great number of stress concentration
points, which are harmful to the mechanical performances, especially when the dispersion and
distribution of GCPs are not uniform. This is probably the main reasons for the degradation of
mechanical performances in many reported polymer nanocomposites [60-65].
1.3.2 Tribological properties
For most polymeric, they have low wear resistance and high frictional coefficient, causing two
main problems: a great mount of wear debris and lowered product life-span. Compared with other
nanomaterials, the graphitic structure in GCNs also contributes to their superior tribological
properties, showing extremely high wear resistance and low frictional coefficient. Thus, the
modification of polymers by GCNs could not only improve properties and functionalities of
polymers, but also significantly improve the tribological properties of polymeric materials, and
consequently prolongs the duration of polymer nanocomposites, which is crucial to both industrial
and commercial applications. The main mechanisms for superior tribological properties of GCN
reinforced polymer nanocomposites include (1) high load-bearing capacity GCNs, reducing the
stresses applied to the polymers; (2) solid lubricator as a result of layered graphitic structures; (3)
High thermal conductivity reducing wear produced by thermal softening; and (4) Reinforced
transfer films by GCNs produced during wear [66-69].
1.3.3 Electrical properties
Polymers are famous for being electrically insulating, with an electrical conductivity as low as
S/cm. With the addition of conductive nanomaterials in insulating polymer matrix, the
resulting composites could exhibit a sharp insulating-conductive transition above a certain loading
level, which is so called percolation phenomenon. This critical loading level is named as
percolation threshold, above which the conductive polymer nanocomposites are typical
semiconductor. However, after percolation, the conductivity usually does not show significant
increase with further increase loading of conductive nanomateirals, corresponding to the “level-off
conductivity”, as shown in Fig. 1.5. The percolation conduction in conductive polymer
nanocomposites can be simply described by equation (1.2)
-0.5 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0
Concentration of GNPs in PEI (wt%)
Figure 1.5 Typical Percolative conduction phenomenon in Polyetherimide (PEI) / Graphite
Nanoplatelet (GNP) composites
c c (1.2)
where σ is the electrical conductivity after percolation occurs, σc is the electrical conductivity of
conductive additives, ϕc is the percolation threshold, and t is a constant solely dependent on the
dimensionality of the additives .
The percolation conduction is strongly related to the formation of conductive network in the
polymer matrix, as a result of contact conduction and tunneling conduction, both of which are
directly determined by the conductivity, geometry, concentration, dispersion and distribution of
conductive additives in polymer matrix. Two common goals dominate the research on conductive
polymer composites: low percolation threshold as well as high level-off conductivity. Especially, in
order to reduce the product cost and keep the superiorities of polymeric materials in low mass
density, flexibility and transparency, the low percolation threshold has been highly pursued.
According to previous reviews on conductive CNT/ polymer, CNF/ polymer and GNP/ polymer
nanocomposites, the high conductivity and aspect ratio of conductive additives usually give the
composites low percolation threshold. Typically, the 1-D GCNs, that is, CNT and CNF, could make
lower percolation threshold than 2-D GNP. At the same time, desirable distribution and dispersion
of conductive fillers (detailed discussion on dispersion and distribution in section 1.), mainly
controlled by processing approaches as well as the interactions between nanofillers and polymers,
are also critical to the resulting electrical properties. One of the lowest percolation threshold found
in epoxy/ CNT composites was as low as 0.0025wt% loading of CNT, leading to an dramatic
increase of electrical conductivity from 10-16
S/cm , however, in another two studies on
epoxy/ CNT nanocomposites, the percolation threshold were, 0.0052vol%  and 0.062 wt% ,
respectively, suggesting the multiple impact factors on electrical properties of polymer
nanocomposites including properties of polymer and GCNs, as well as composite processing.
Some other reported low percolation thresholds include: 0.05 vol% in polypropylene /multi-walled
CNT nanocomposites prepared by shear mixing ; 0.011vol% in polystyrene /GNP
nanocomposites prepared by in-situ polymerization ; 0.9 vol% in silicone rubber /GNP
nanocomposites prepared by wet mixing and curing , etc.
Regarding the level-off conductivity of conductive polymer nancomposites, it usually stops at
S/cm with increasing GCN loading, at a reasonably low loading level without
compromising properties of polymeric materials, compared with high electrical conductivity of
GCNs as shown in Table 1. Besides choosing GCPs with higher conductivity, in order to realize
high level-off conductivity at a low loading level, some efforts have been made, such as dispersing
aligned, straight and unentangled CNT in epoxy via a shear intensive mechanical stirring process
followed by the application of low shear force to reagglomerate CNTs. In resultant nanocomposites,
the maximum conductivity can reach ca. 10~2
S/cm at 1wt% CNT loading . Besides the
application of aligned and /or oriented GCNs, another useful approach to increase maximum
conductivity without significantly increasing the GCNs loading is the use of hybrid nanofillers, for
example, the application of (CNT + carbon black (CB)), as well as (GNP+ CB) and
(GNP+CNT+CB) hybrid systems in epoxy [77,78]. In these studies, the hybrid nanofillers have
successfully improved the electrical conductivity, compared with single nanofiller reinforced
polymer nanocomposites, with lower percolation threshold and higher level-off conductivity by
controlling the ratio of different carbon nanomaterials.
Figure 1.6 Different surface resistivity levels and their applications (provide by Sabic Innovative
Most research has been only discussing volume electrical properties, and ignored surface
electrical property which is the measurement of the resistance of materials to the flow of electrical
current across their surfaces. Usually, it does not show exactly the same percolative conduction
behavior as volume conductivity does . Surface conductivity is a measurement of the ability of
materials surface to transport charger carriers, which is considered as a critical factor to antistatic
coating/materials which have important applications in aerospace/ground transportation and
electronic packaging, etc. Fig. 1.6 presents the requirement of surface resistivity in different
applications. For a typical antistatic application, the surface resistivity below 1012
ohm/sq is desired.
While for a faster static dissipation rate, it is expected to lower than ~109
ohm/sq. According to our
studies , the improvement of surface electrical properties was more difficult than volume
electrical properties, which needed more loading of conductive GCNs. Thus, the effective decrease
of surface resistivity of polymer nanocomposites at a low loading level is a rather challenging
research subject. However, the study on surface electrical properties is still insufficient.
1.3.4 Dielectric properties
Low dielectric constant is an important characteristic of polymers, and is usually between 2~10.
This makes polymers ideal materials for high speed integrated circuit applications. However, with
the increasing demand of light weight, flexible and/or transparent electronics, sensors, capacitors as
well as battery, the modification of dielectric properties of polymers are necessary.
Dielectric property is a result of short range conduction of materials, unlike electrical
conductivity which represents long range conduction behavior. Dielectric constant, or relative
permittivity (ε’), is a measurement of material ability to polarize in electrical field (formation,
orientation and alignment of dipoles), as a result of various polarization mechanisms including
electronic, atomic and ionic polarization. In heterogeneous materials systems, the polarization
usually occurs at interface between different phases/materials, due to the accumulation of charges at
the interface. The interfacial polarization, so-called Maxwell-Wagner-Sillars polarization, is
considered the main polarization mechanism in polymer nanocomposites . Basically, the
dielectric properties as a result of interfacial polarization in a two-phase dilute random suspension
of spherical particles can be described by Maxwell-Wager equation (Eq. 1.3):
* ** *
* 2 * * 2 *
m p m
where, ε*, ε*m, ε*p are the complex permittivity of suspension, matrix as well as particles,
respectively; ϕ is the volume fraction of particles. In particular, the complex permittivity is given
by equation 1.4:
* ' "i (1.4)
where, ε’ is the relative permittivity, i.e. dielectric constant; ε” is the dielectric loss resulting from
the energy dissipation during polarizations (alignment and orientation of dipoles). Depending on
specific applications, both low and high ε’ are desired. For example, for the application in high
speed integrated circuit, low ε’ is required in order to obtain high efficiency; while the energy
storage applications, such as capacitors and batteries, high ε’ is needed. GCNs can significantly
improve the dielectric constant of polymer nanocomposites at low loading levels, compared with
ceramics and metals. For example, with addition of only 2.34vol% GNPs, the dielectric constant of
Polyvinylidene Fluoride (PVDF) can be dramatically elevated from ~8 to ~107
at 1000Hz .
However, because of the complex polarization mechanisms, which do not only depend on the
concentration, dispersion and distribution of GCNs, but also strongly relate to chemical structures
and dielectric constant of constituent materials in the nanocomposites, the dielectric response of
polymer nanocomposites is a more complicated phenomenon. This can be reflected in a similar
study on PVDF/GNP composite, in which the dielectric constant only reached as ~102 at 1000 Hz at
the same loading level .
Dielectric loss yields during the formation and orientation of dipoles (polarization). In the
pursuit of high dielectric constant, low dielectric loss is also expected for high dielectric efficiency
and long service life of dielectrics. However, high dielectric constant usually leads to high dielectric
loss, as a result of intense polarizations. Compared with increasing dielectric constant, keeping the
dielectric loss at a low level (close to or lower than polymer) seems a much more challenging task.
Few current attempts for this purpose include (1) introducing insulating layers between
nanoparticles and polymer matrix, such as epoxy/nano-Ag composites , and (2) fabricating
sandwich structure polymer nanocomposites, such as PVDF/CNF nanocomposites .
1.3.5 Damping properties
High performance polymeric materials possess high specific strength and specific modulus, i.e.
high strength and modulus to density ratio, which leads to their increasing popularity in construction,
ground transportation and aeroplane. However, these polymers usually exhibit poor ability to damp
mechanical vibration and reduce noise, due to rigid polymer chain structures and weak inter-/intra-
molecular frictional motion which cause high damping capability of low modulus elastomers. Thus,
the low damping capability of high performance polymers blocks their rapid development in these
To improve damping capability of high performance polymeric materials without sacrificing
their outstanding mechanical properties is one of the most challenging research areas in polymer
nanocomposites, since high damping capability and high strength/modulus are usually considered as
two inversely related aspects of material properties. This is probably the main reason for very few
successful attempts reported on this topic. A commonly used strategy for this purpose is to reinforce
polymers by rigid fillers, such as glass fiber, carbon fiber , and CNTs, and improve frictional
motions among fillers and polymers. In particular, CNTs proved to be the most successful in
enhancing damping capability [83-85]. However, simply adding rigid fillers to a polymer matrix
could only increase the stiffness and reduce the damping capability , due to restrained polymer
chain mobility and insufficient frictional motion. Accordingly, instead of direct mixing rigid fillers
with polymers, continuous CNTs  as well as CNT thin films bonded by epoxy adhesive have
been introduced in polymer nanocomposites to achieve high damping properties, while maintaining
other mechanical properties .
1.3.6 Other properties
Thermal stabilities: Another well-know limitation of polymeric materials is their low thermal
stability compared with metals and ceramics, which consequently restricts the high temperature
applications of polymers. Thermal stability includes two aspect chemical stability evaluated by
onset temperature (To) of thermal degradation/decomposition, and physical stability evaluated by
the coefficient of thermal expansion (CTE). Thermally stable GCNs have proved to be efficient
functional fillers in increasing To of polymers. Several mechanisms have been proposed to explain
the contribution of GNPs to the improved the thermal stability, here, we list some of them: 1) the
homogeneously dispersed GCNs acts as “efficient heat sinks”, which consumed more heat than the
matrix and did not allow the accumulation of heat within the latter, and thereby prevented oxidation
at the early stages of degradation ; 2) the homogeneously dispersed GCNs could serve as the
mass transfer barriers (shielding effect) against the volatile pyrolized products ; 3) the
interfacial polymer phases in the vicinity of GCN surfaces are restricted by the bonding from GCPs,
and the energy needed to decomposition would increase, alter the ability of degraded molecules to
diffuse and evaporate.
Thermal expansion is very common to polymeric materials at high temperature. The addition of
rigid fillers, such as GCNs  to polymer matrix is a useful way to restrain the thermal expansion,
due to the confinement effect of rigid fillers with low thermal expansion. CTE usually increases
with temperature, and decreases with increasing loading and uniform dispersion. In our previous
study, by applying a novel two-step melt processing methods to polycarbonate/ CNF composites,
the CTE of nanocomposites showed a 30% decrease at 1wt% CNF loading, compared with the
nanocomposites fabricated by conventional one-step processing method, due to significantly
improved dispersion quality . Compared with other 0D and 1D nanofillers, e.g. CB, CNF and
CNT, the excellence of GNPs in improving the dimension stability is obvious, according to studies
on polypropylene nanocomposites reinforced by various carbon nanomaterials .
Thermal conductivity: Low thermal conductivity (low heat dissipation capability) limits the
applications of polymers in liquid cooling and ventilation garment, power electronics, electric
motors and generators, heat exchangers, etc , despite their light weight, corrosion resistance,
lower manufacturing cost and ease of processing. Due to the high thermal conductivity of GCNs,
the incorporation of GNPs in polymer matrix usually leads to the increase of thermal conductivity
of polymers [92-93]. Thermal conductivity is primarily determined by the vibration of lattice
phonons, and also by thermal motion of electrons , which relate to dispersion, orientation and
aspect ratio of GCNs, as well as interface and contact thermal resistance in polymer nanocomposites.
Various models have been developed to predict the thermal conductivity of polymer
nanocomposites , however, owing to the insufficient understanding of thermal conduction
mechanisms, especially, lattice vibration behavior in polymer nanocomposite, the discrepancy
between experimental and theoretical values is still remarkable.
The GCNs with high aspect ratio lead to better thermal conductivity , owing to a smaller
contact resistance . The study of thermal conductivity of Epoxy/GNP nanocomposites
suggested the importance of interface thermal resistance on the thermal conductivity of composites.
The same results were also found in CNT reinforced polymer nanocomposites. After the GNPs were
treated by nitric acid, the interfacial bonding was improved due to the existence of polar groups on
the GNPs, so is the interface thermal resistance. Consequently, the thermal conductivity of nitric
acid treated GNP has better modification efficiency . In Ethylene-co-Vinyl Acetate
(EVA)/graphite oxide (GO) nanocomposites, the thermal conductivity decreased at highest loading
(4phr), as a result of aggregation of fillers . Compared with CNF and MWCNT, the GNP filled
PEI nanocomposites have the highest thermal conductivity along direction of nanofiller alignment
in a wide loading range, no matter solution or melt mixing was applied to prepare the composites
. In the same study, the unoriented composites have also been prepared, still discovering the
highest thermal conductivity in GNP /PEI nanocomposites.
Figure 1.7 Morphology of as-received CNF in aggregate form
1.4 Dispersion and distribution of GCNs in polymers
It is well recognized that in polymer nanocomposites, the dispersion and distribution of
nanomaterials are of greatest importance to the properties and functionalities of polymer
nanocomposites. In majority of applications, a uniform dispersion is required, such as mechanical,
electrical and dielectric properties. However, due to the extremely high surface energy,
nanoparticles tend to aggregate in nature. For 1-D nanoparticles, such as CNF and CNT, the severe
entanglement also exists, which makes the dispersion more difficult, as shown in Fig.1.7.
Figure 1.8 Schematic sketches of the effect of 1D GCNs on the conductivity of polymer
In fabricating high performance functional polymer nanocomposites according to specific needs,
the effectual control of dispersion and distribution is the most fundamental and challenging task,
and have been the primary research subjects and goals in developing polymer nanocomposites. Fig.
1.8 summarized four typical dispersion and distribution states that frequently occurred in
polymer/1-D GCNs nanocomposites, revealing the direct relationships between electrical
conductivity and dispersion/distributions of 1D GCNs. Current technologies for controlling
dispersion and distribution can be divided into two categories: optimum composite processing and
surface modification of GCNs.
1.4.1 Dispersion and distribution control during composite Processing
Polymeric materials can be processed into desired forms in both melt solution states. In
particular, melt processing (extrusion, injection molding and melt mixer, etc.) is more industry
favorable due to its high product yield and consistence, in which, the high melt shear force can be
transferred to the GCN agglomerates and results in disentanglement /separation of GCNs in
polymer melt. However, melt shear force is usually insufficient to uniformly disperse GCNs in
polymer melts, especially for CNTs with higher aspect ratio and surface energy. Thus,
ultrasonication is frequently used, either before or during melt processing. Ultrasound could
provide high energy to break down the GCN agglomerates. Previous study on CNF revealed that
ultrasonication at a certain power level could remarkably shorten CNFs, leading to the
disentanglement and uniform dispersion , thus ultrasonication is a key step in pre-treatment of
GCNs and solution processing of polymer nanocomposites. In order to improve the dispersion
quality of nanomaterials, various advanced processing techniques have been developed. Zhong et al
 developed a lean two–step processing technique using solid “nano-nectar” assisted melting
dispersion. Ultrasonication has also been directly applied to polymer melt during extrusion 
by introducing an ultrasound supplier to a traditional extruder. For some shear-sensitive polymers,
high shear force will cause degradation or other structure changes during melting processing. In
order to disperse nanomaterials in this group of polymers, Yang et al  directly injected uniform
nano-suspensions prepared by ultrasoncation into extruder during mild melt compounding process
with low melt shearing.
1.4.2 Improvement of dispersion via surface modification
As mentioned earlier, the challenging nanoparticle dispersion in polymer matrix is due to their
high surface energy which is determined by their surface properties. Thus, in addition to composite
processing, surface modification is another important and efficient approach to realize uniform
dispersion. In general, surface modification has two functions: decrease surface energy and improve
interfacial bonding between polymer and nanomaterials. For GCNs, they are composed of SP2
hybridized carbons which are chemically stable and incompatible with polymers. The surface
modification of GCNs includes covalent and nano-covalent approaches.
Covalent surface modification of GCNs involves two steps: oxidization of GCNs and
modification by surfactant. Because of the chemically stable SP2 hybridized carbons, the traditional
surfactants used in modifying inorganic mineral fillers, such as CaCO3 and clays, with active
surface groups, cannot be attached to the surface of GCNs. Thus, before any reactions with
surfactants, it is necessary to create reactive sites on the GCN surfaces. Typically, this has been
done by various oxidization procedures, such as mixture of concentrated H2SO4/HNO3/KMnO4,
ozone and radiations. Consequently, a great amount of functional hydroxyl and carbonyl groups are
generated on the surfaces and edges of GCNs, providing reactive sites for further modification.
After the oxidization, the surfactant, such as silane coupling agent, could react with hydroxyl and
carbonyl groups to form strong and stable covalent bonding with GCNs, while the other end of
surfactant molecules usually forms strong interactions with polymer chains, leading to strong
interfacial bonding between GCNs and polymers. Covalent surface modification has proved great
excellence in improving mechanical and tribological properties in polymer nanocomposites [42, 57].
In recent years, non-covalent surface modification of GCNs has been gaining more and more
popularity. Non-covalent surface modification bonds GCN surfaces and surfactants via π-π stacking,
thus, the surfactant must contain conjugative structures which commonly exist in conductive
polymers, such as polythiophene, polypyrrole and polyaniline, etc. Regarding the compatibility, the
copolymers of conductive polymers are frequently used for non-covalent modification, for example,
poly(3-hexylthiophene-co-styrene) in polystyrene (PS) /CNT nanocomposites . The
polythiophene block in the copolymer could attach to the SP2 hybridized CNTs, while the PS block
is compatible with PS matrix, accounting for the good interfacial bonding as well as uniform
dispersion of CNTs.
In addition to the contribution of non-covalent surface modification to uniform dispersion of
GCNs, it is also worthy to mention its advantages in developing functional polymer nanocmposites,
compared with covalent surface modification.
First of all, non-covalent surface modification could substantially retain extraordinary electronic
properties of GCNs. It is well known that the uniqueness of GCNs comes from the SP2 hybridized
carbons. However, the oxidization in covalent surface modification could convert them into SP3
hybridized carbons which devastate both mechanical and physical properties. This situation could
be completely avoided in non-covalent modification, because of needlessness of oxidization in non-
covalent modification. Secondly, conductive polymers have proved their potential in many
important applications including optics, electronics and energy, etc . The non-covalent surface
modification of GCNs with unique electronic structures could further improve and explore
magnificent functionalities by integrating excellences of conductive polymers and GCNs. Further,
the non-covalent modification provides more exciting opportunities for engineering super-
molecular surface structures for advanced applications .
1.5 Problem statement
PEI has shown great potential in a variety of advanced applications. In particular, due to its
superior mechanical and thermal stability at elevated temperature, it provides great opportunity to
develop functional polymeric materials for high temperature applications, such as high temperature
capacitors and batteries. In addition, considering the superior mechanical, thermal and electronic
properties of GCNs, it is expected that the nano-reinforcement of polyetherimide by GCNs will be
one of the most efficient ways to realize high performances and functionalities of PEI matrix
However, despite the fascinating potential of PEI/GCN composites, it is still an insufficiently
studied research area. For example, according to ISI web of knowledge, the number of research
papers on PEI/GCN composites was only 58 using PEI, carbon and graphite as title keywords,
(based on search results on April 11th
, 2012), in contrast to 355 papers on polyimide /GCN
composites. Thus, the studies on PEI/GCN nanocomposites as well as their structure-property
relationships are necessary and indeed urgent.
In particular, as demonstrated in previous sections, dispersion and distribution of nanomaterials
are critical to the properties and functionalities of polymer nanocomposites. Although PEI has a
better processibility, as a result of flexible ether groups, compared with other aromatic PI materials,
its viscosity is still high enough to make dispersion and distribution of nanomaterials very difficult,
which prevents the commercialization of PEI/GCN nanocomposites. Also, the studies on PEI/GCN
composites usually only focused on achieving uniform dispersion of GCNs to modify mechanical,
thermal and electrical properties, etc. [92, 105], while overlooking the possibility to obtain high
performance and new functionalties by varying dispersion and distribution of GCNs in PEI matrix.
In other words, the relationship between variability of GCN dispersion/distribution and
functionalities of polymer nanocomposites is still a rarely exploited research area, the study of
which is believed will cultivate new and fascinating properties and functionalities and boost the
development of advanced polymeric materials. Thus, besides achieving uniform dispersion, the
effectual control of dispersion and distribution of GCNs in PEI is also a substantial task to the
potential of PEI/GCN nanocomposites, to satisfy applications requiring extraordinary mechanical
and thermal stability as well as low mass density.
In addition to controlling dispersion and distribution of GCNs, surface modification of GCNs is
another useful approach to realize property improvement and new functionalities. The selection of
proper surfactants as well as modification procedures is a key factor in fabrication of high
performance polymer nanocomposites. Not only the structures and properties of polymer matrices,
GCNs and surfactants, but also the interaction among them should be comprehensively considered,
in order to obtain desired properties and/or functionalities. Although a great number of works on
surface modified GCNs reinforced polymer nanocomposites have been reported, the understanding
of the effects of surface modification on structure and properties of polymer nanocomposites is still
insufficient, because of complicated physical and chemical phenomena determined by materials
structures and properties as well as composite processing techniques. Thus, in order to develop
multifunctional PEI/GCN nanocomposites, the surface modification of GCNs deserves more
attention and effort.
PEI is considered as an ideal material for intra-structure for airplane and ground transportation,
due to its low mass density and excellent mechanical and thermal properties. However, as a typical
insulator, its low static dissipation rate might cause accident as a result of electric discharging
phenomenon. Thus, to improve the static dissipation capability of PEI is important. Undoubtedly,
the addition of conductive GCNs could significantly improve the electrical conductivity of PEI
composites at a low loading level. A comprehensive understanding of static dissipation capability
of polymeric materials should include three aspects: volume electrical conductivity, surface
electrical conductivity as well as dielectric constant, among which, surface electrical conductivity
and dielectric constant are usually neglected in most reported works. The fact is surface electrical
conductivity is as important as volume electrical conductivity, because the accumulated electrical
charges are mostly dissipated through materials. Moreover, dielectric constant suggests the ability
of materials to store electrical charges; as a consequence, a high dielectric constant slows the static
dissipation rate. The reinforcement of polymers by GCNs could not only increase electrical
conductivities, but also increase the dielectric constant of resultant composites, as demonstrated in
Section 1.3.3 and 1.3.4. Thus, it is necessary to evaluate both electrical conductivities and dielectric
constant of PEI/GCN composites to truly understand their static dissipation capability and provide
constructive information for optimizing material structures and achieving satisfactory static
1.6 Research objectives and significances
As the core part of Boeing/NSF projects, this project aims at understanding properties and
functionalities of aeropolymer polyetherimide and developing high performance nanocomposites
reinforced by graphitic carbon nanofillers (GCNs). The main objectives include:
1. Effectual control of GCN dispersion and distribution in PEI matrix, and correlating with
functionalities and properties. As demonstrated earlier, the dispersion and distribution of
GCNs are the most critical factors affecting the properties and functionalities of polymer
nanocomposites. A comprehensive understanding of their relationship is important to
developing advanced nanocomposites. In this study, different polymer processing techniques as
well as multiple GCNs treatment approaches were applied to obtain different
dispersion/distribution states. Furthermore, the successful establishment of functionality –
dispersion/ distribution relationship is also meaningful to non-destructive evaluation (NDE) of
quality of polymer nanocomposites. With increasing needs of polymer nancomposites in many
siginicant applications, such as aerospace/airplane, ground transportation and electronics, etc,
the quality control of polymer nanocomposite becomes more and more important, in particular,
the NDE of dispersion/distribution of GCNs in a polymer matrix is the center of quality control
of the polymer nancomposite, which is highly demanded in nanocomposite industry.
2. Understanding effects of surface modification of GCNs on properties and functionalities of
polymer nanocomposites. Obviously, the industry is calling for green manufacturing
techniques, considering the environmental issues as well as workers’ health concerns. One of
my research goals is to realize functionalities of polymer nanocomposites by controlling
dispersion and distribution of GCNs, as well as structuring polymer nanocomposites, while
avoiding the use of toxic and volatile organic chemicals and complicated chemical procedures.
However, numerous studies [57, 103, 105] also revealed that proper surface modification of
GCNs could not only further improve properties of polymer nancompsoities, but also
substantially dig out new functionalities of GCNs and their polymer nanocomposites. Thus,
according to specific applications, both covalent and non-covalent surface modifications were
carried out on GCNs, and their effects on the properties of resultant polymer nanococmposites
were studied, referring to pristine GCNs.
3. Structuring bulk composites for high performances and new functionalities. In addition to
controlling dispersion/distribution of GCNs in polymer matrix, as well as surface modification,
structuring bulk composite is also a useful approach to improve properties and obtain new
functionalities of polymer nanocomposites. In particular, multilayered or layer-by-layer
composites have proved their excellences , compared with traditional mono-layer
composites. In this research, bi-layer nanocomposites have also been developed in order to
further improve the properties of polymer nanocomposites and explore their new functionalities.
As one of the most important engineering polymers, polyetherimide is showing great potential
in advanced applications, especially high temperature applications, due to its excellent mechanical
and thermal stability. Thus, developing new high performance and multifunctional materials based
on polyetherimide is intriguing and valuable. At the same time, compared with other nanomaterials,
GCNs prove to be the best in comprehensively improving fundamental properties (mechanical,
triobological, thermal properties and so on) and realizing multifunctionalities (optical, electrical,
dielectric properties and so on) of polymeric materials. Therefore, it is promising to combine the
superiorities of GCNs and PEI to fabricate advanced polymer nanocomposites with outstanding
mechanical and physical properties for a variety of applications from commodity, industry to
military and aerospace, etc. Furthermore, this research also focuses on in-depth understanding of
structure-property relationships in nanocomposites, which could provide more insightful and useful
knowledge for both academic research and industrial production of polymer nanocomposites.
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Single Negative Metamaterials in Unstructured Polymer Nanocomposites
toward Selectable and Controllable Negative Permittivity
Organic nanocomposites fabricated in our lab exhibited negative permittivity, a characteristic
heretofore reported only in artificial metallic metamaterials. Metamaterials are a class of material-
like structures with negative permittivity and negative permeability which have attracted great
attention because of their potential to realize particular electromagnetic/optical properties that do
not exist in naturally occurring materials. It has been recognized that the unusual properties of
metamaterials are a consequence of their specially designed structure, and are independent of their
compositions. However, our recent studies have shown that negative permittivity can be attained
from non-metallic nanocomposites without specially designed structural attributes. Additionally, in
contrast to reported research, the organic polymers were found to be influential to the establishment
of pronounced negative permittivity. Moreover, these negative permittivity levels were strongly
dependent on the composition factors of the nanocomposites, such as the concentration levels,
lengths and micro-structures of the nano-fillers, as well as the chemical structures of polymer matrix
materials. It is believed that such foundational knowledge can lead to greater understanding of the
prerequisites for achieving organic metamaterial films with negative permittivity and negative
permeability. In many applications such nanodielectric materials would enable greater efficiency in
structural and electronic systems, significantly benefiting band gap, waveguide, wave filter and
Negative permittivity is one of the key characteristics of the unique metamaterials, a novel class
of artificial materials with special structures which simultaneously exhibit negative permittivity and
negative permeability. In metamaterials, adherence to the right-hand rule, which is obligatory for all
other materials working in an electro-magnetic field (electric field E, magnetic field H), does not
apply. The wave vector k in metamaterials is consistent with a left-handed rule, and, hence, this
special class of materials is also called left-handed materials (LHM), first theoretically proposed by
Veselago in 1960s 
. The negative permittivity and negative permeability endow LHM with
several special properties, such as negative refractive index, reversed Doppler Effect, and reversed
Cherenkov radiation 
, none of which exist in natural materials. Metamaterials can be applied in
sub-wavelength imaging, cloaking, wave filter, super lens and band gap design [3-5]
. Since 2000,
when investigations of metamaterials were experimentally performed by double slit ring resonator
and linear vibrator 
, research on the construction and application of metamaterials in optical and
microwave frequencies range have attracted great interest, and much theoretical work has been
published, however the realization of their promise still requires great effort.
It is well recognized that the unusual properties of metamaterials are determined by their
specially designed structures, rather than their compositions. In other words, the permittivity and/or
permeability are not, or are just weakly dependent on chemical structures and constituent
compositions, or interactions among different components in the materials. Various structures have
been designed by physicists and electrical engineers to obtain negative permittivity and negative
permeability, including double split ring resonators 
, S-shaped resonator 
, multilevel dendritic
, nano/small aperture [8-9]
, metallic nanoclusters
, among others. In all of these,
resonance was considered to be the main mechanism of the realization of both negative permittivity
and negative permeability, while the phase changes close to resonant frequency 
. In contrast, the
exploration of metamaterials from the point of view of materials (chemistry and properties) has not
attracted substantive attention from materials researchers. Among the very limited reported work,
almost all research in this field focused on metals and ceramics which contain metallic elements and
can consequentially realize negative permittivity and/or permeability under certain conditions. In
particular, ferromagnetic materials were the subject of considerable interests in realization of
metamaterials, owing to the resultant negative permittivity of metallic materials below plasma
frequency as well as negative permeability of some ferromagnetic metals and metallic alloys in the
vicinity of the ferromagnetic resonance frequency[11-13]
. However, to our knowledge, research has
been rarely reported on negative permittivity or negative permeability directly obtained from
conventionally fabricated materials .
On the other hand, truly practical applications aspire to large scale, light weight, deformable and
efficiently processable metamaterials which are not fully realized by those costly and precisely
designed structures, or metallic and ceramic materials. Based on recent success and popularity of
polymeric materials in fabricating large area and flexible electronic devices 
, it is compelling to
consider the fabrication of polymeric metamaterials with the advantages mentioned above. For the
polymers involved in a metamaterial system, it was usually considered that they were only applied
as insulating host/substrate in metamaterials [11,12,16,17]
. Moreover, their negative permittivity and/or
negative permeability were considered to be caused solely by those inorganic inclusions and were
independent of the polymer properties. In a recent study, on nanopolyanilne/epoxy hybrids obtained
by a special absorption-transferring process
, the occurrence of negative permittivity was
considered to have resulted from the formation of a continuous conductive network of conductive
2.2.1 Raw materials
Polyetherimide (ULTEM 1000) and Polyimide (EXTEM XH1015) were supplied by Sabic
Innovative Plastic Inc. The relative permittivity of their cast coated films is 1.7 and 2.1, respectively,
measured at 1MHz in our study. Cup-stacked carbon nanofibers (Pyrograf® III), with diameter of
60nm-150nm and length of 30μm-100μm, were purchased from Applied Science Inc. As-prepared
herringbone carbon nanofibers, with the diameter of 50-150nm and length of 50 µm-100µm were
supplied by Vanderbilt University.
2.2.2 Preparation of polymer nanocomposites
The ultrasonic treatment of CNFs was processed on a Branson digital ultrasonicator (Model 450)
at 20% amplitude for 1 hour in the presence of the diluent Acetone. Each time, only 0.3g CNFs
were treated and the mass/volume ratio of CNFs and Acetone was fixed at about 0.3g/12ml. After
the treatment, all the CNFs were dried at 700C, at the same time, maintaining stirring to prevent
agglomerating the CNFs during drying. Thus, the shortened CNFs were obtained and placed on
After drying at 100°C for about 2 hours, the as-received and ultrasonically treated CNFs were
incorporated into the PEI(PI)/Dichloromethane solution with a mass/volume ratio of 1g/3ml,
respectively. The concentration of CNFs is given in Fig.1. The combined materials were then
preliminarily mixed on a spin mixer for about 12 hours. The spin mixer does not offer a sufficiently
strong shear force to implement homogeneous dispersion of CNFs in polymer matrix, but it is
effective to assure the penetration of polymer chains into CNFs network. After that, the suspension
was immersed in an ultrasonic cleaner (Branson 1510) for consecutive two and half hours to expel
the bubbles in the suspension. The temperature of water bath was kept between 250C and 30
The films with thickness of about 0.06mm were cast coated on a glass plate. After the solvent
completely evaporated, the coated glass plate was immersed in cold water for a while in order to
separate the films from the glass plate.
2.2.3 Morphology analysis
The microstructures of PEI/CS-CNF (as-received) and PEI/CS-CNF (ultrasonically treated)
nanocomposites were analyzed by Optical Light Microscope (Olympus BX51) at the magnitude of
×5 and Scanning Electron Microscope (SEM, FEI 200F) at the magnitude of ×20000. Before SEM
observation, the surface of film samples was rinsed and etched by Dichloromethane for the
convenience of observing morphologies of CNFs in PEI matrix.
2.2.4 Dielectric and electrical tests
The dielectric tests in AC field were conducted on a Universal Dielectric Spectrometer BDS 20
from Novocontrol Inc. The input voltage (Vrms) was 1V, and the test frequency range was set
between 100Hz and 3MHz. Before the tests, film samples were rinsed by Acetone to remove
possible impurities on film surfaces to reduce their interferences. Only the real part (relative
permittivity) of the complex permittivity was discussed in our study. For each polymer
nanocomposites in our study, at least 15 film samples were tested.
2.3 Results and discussion
Our latest studies of polymer nanocomposites prepared in our facilities surprisingly showed
negative permittivity values. It has been discovered that the negative permittivity can be obtained
from lab-fabricated non-metallic nanocomposites by conventional solution processing method,
without any specially designed structures within the nanocomposites, additionally, in contrast to
reported research [11-12]
, the organic polymers were found to be influential to the establishment of
pronounced negative permittivity. The negative permittivity levels are clearly dependent on the
composition factors of the nanocomposites, such as the concentration levels, aspect ratio and micro-
structures of the nano-fillers, as well as the chemical structures of polymer matrix materials,
contrary to the above beliefs.
2.3.1 Physical origins and concentration dependence of negative permittivity
Figure 2.1 Frequency dependence of permittivity of PEI/CS-CNFs nanocomposites: Except for the
nanocomposites containing 1.0wt% CNFs, all other PEI/CNF nanocomposites show negative
permittivity in the frequency range of 1KHz-3MHz, while the PEI/1.0wt%CNFs nanocomposite
only shows negative permittivity in the vicinity of resonance (2KHz- 10KHz). Obviously, the
increasing loading of CNFs boosts the resonance phenomenon around 5KHz. However, this
resonance frequency is almost independent of the composition of PEI nanocomposites.
In the nanocomposites there is no metal wire involved; rather the fillers primarily used are
commercialized carbon nanofibers (CNFs), which are extensively applied in polymer
nanocomposites for improving mechanical, electrical, dielectric and thermal properties of polymers.
The CNFs have distinctive electric and dielectric properties different from metallic wires: the
conductivity of non-metallic CNFs (usually lower than 103 S/m) is much lower than that of metallic
wires (about 107
S/m), furthermore, owing to the graphitic nature, the permittivity of CNFs is
usually positive and should be approximately equal to that of graphite which is in the range of 12-15,
unlike metallic materials whose permittivity manifestly becomes negative below plasmon frequency
(usually in ultraviolet region) 
. In our work, highly purified cup-stacked CNFs (CS-CNFs,
Pyrograf® III from Applied Sci. Inc, iron catalyst content <100ppm), characterized by a stack of
truncated conical graphene layers (cups) held by Van der Waals forces, were primarily utilized. For
matrix material, polyetherimide (PEI, Ultem 1000, Sabic Innovative Plastics Inc.), a derivative of
polyimide (PI) by introducing flexible phenylene ether groups to PI chains, was applied as the
polymer matrix to fabricate the CNF nanocomposites.
Initially, we obtained negative permittivity from CNF/PEI nanocomposites. Moreover, it was
discovered that the concentration levels in the nanocomposite were shown to exhibit significant
influence on the negative permittivity values.
Negative permittivity was obtained in the frequency range of 0.1 KHz - 3 MHz in PEI / CS-
CNFs nanocomposites films. It should be noted that, deviating from the widely accepted opinion of
structure-determined metamaterials proposed by Roger M. Walser 
, the negative permittivity
shows remarkable composition dependence: the absolute values of negative permittivity were
enhanced by increasing CNF loading, as shown in Fig. 2.1. A resonance phenomenon appeared
around 5 KHz at which the negative permittivity was the highest, and the composition dependence
of negative permittivity was also the strongest. It is believed that this resonance is the causality of
the negative permittivity in our nanocomposites film samples. According to the J.B. Pendry’s model
[17, 19], a 3D network of metallic thin wires could exhibit negative permittivity below GHz frequency
(low frequency plasma), as long as the wires are thin enough (radius of 1.0×10-6
m in Ref. ) and
Figure 2.2 Optical microscope photos (×50) and SEM images (×20,000) of PEI nanocomposite
films: (A). PEI/CS-CNF (2 wt % as-received) nanocomposites. According to the optical image,
there are plenty of agglomerates (black regions) and a few separately distributed long CNFs. The
SEM image shows that CNFs are intensely entangled in the agglomerates. (B). PEI/CS-CNF (2wt%
ultrasonically treated) nanocomposite. After the ultrasonic treatment, the bigger size agglomerates
break down via shortening of as-received CNFs. Thus, the dispersion of CNFs is greatly improved.
Separately dispersed long CNFs
Improved dispersion of
ultrasonically treated CNFs
have good continuity (long wires) - the maximum diameter of CNFs in our study is about 150 nm
and satisfies the requirement of thin wire. Also, in J.B. Pendry’s model, an insulating polymer
substrate was applied in the construction of 3D network to adjust the density of wires and
consequently plasma frequency. An important aspect of this model is the continuity of 3D network.
Figure 2.3 Influences of discontinuity of CNFs networks on permittivity of PEI/CS-CNFs
nanocomposites. After the continuity of the CNF network was disrupted in PEI nanocomposites by
ultrasonically cutting, the permittivity becomes positive. At the same time, the permittivity kept
constant at 1.8 in the whole frequency range, and resonance phenomenon completely disappeared.
The essence of discontinuity is shortening of wires and the disconnection among wires in neighbor
cells. After cutting away 40% of each wire, plasma-like behavior disappeared, and consequently,
the negative permittivity vanished. In our work we studied the effect of continuity of 3D network
of CS-CNFs by ultrasonic cutting. The effect of ultrasonic treatment on cutting long CNFs,
breaking big CNF agglomerates and improving dispersion of single CNF in polymer matrix has
been shown in our previous work 
, and is also seen from the optical and SEM images in Fig. 2.2
(A, B). The results of dielectric testing of PEI/ CS-CNFs (ultrasonically treated) nanocomposite
shows that the permittivity is totally positive in the whole frequency range in our study (Fig. 2.3)
along with the disappearance of resonance phenomenon around 5KHz. This is exactly consistent
with J.B. Pendry’s model. On the other hand, although continuity of the local CNF networks was
destroyed by ultrasonic treatment, the overall continuity of conductive network was improved. This
indicates that the conductivity of PEI/CNF nanocomposites can be considered to be unrelated to the
negative permittivity, which is inconsistent with the results of the nanopolyaniline/epoxy hybrids
According to the microstructures of the film samples of both PEI nanocomposites (Fig. 2.2 (A,
B)), the changes of structures and dielectric properties of the PEI/CS-CNF nanocomposite films
with as-received CNFs and ultrasonication-cut CNFs are illustrated in Fig. 2.4. These industrial
CNFs are in the form of large agglomerates (Fig. 2.2A), owing to high surface energy and
entanglement of the long fibers. It is reasonable to consider that the CNF agglomerate is an
analogue of the metallic 3-D network since the diameter of CNFs is usually in the range of 50nm-
100nm, while they have very high aspect ratio with lengths up to 200microns, effectively providing
continuous 3-D networks of CNFs. The resonance phenomenon in the PEI/CS-CNF nanocomposite
with negative permittivity can be seen to derive from the structural unit (orange color region) in Fig.
2.4. However, this resonance frequency, independent of concentration of the CS-CNFs, is not a
typical plasmon resonance frequency which is usually in ultraviolet region and can be calculated by
ln( / )
n e cw
m a a r
where, wp is the plasmon frequency, and r is the radius of thin wires. The detailed information of
parameters in the equation can be found in Ref. . Moreover, if we substitute wp=5 KHz and
r=100 nm into Eq. (2.1), the calculated lattice constant a is much higher than the dimensions of our
test samples (diameter = 20um), meanwhile, from this equation the plasmon frequency can be
adjusted by the density of wires which would result in the changes of lattice constant (a). Therefore,
the negative permittivity of PEI/CS-CNF nanocomposites is likely not the result of conventional
Figure 2.4 Simulations of structural changes in CNF network caused by ultrasonic cutting and
influences on electric field. Before treatment, when an external electric field is applied, because of
the negative permittivity, the electric field cannot go through the nanocomposites films, and the
electromagnetic waves could not transmit through the films; after treatment, the destroyed
continuity of CNFs network accounts for positive permittivity, so that the applied electric field is
permitted to penetrate the films.
plasmon resonance. In addition, it is also easily understood why the “higher” negative permittivity
occurs at higher loading of CS-CNFs. The incremental concentration of CS-CNFs (as-received)
resulted in a greater number of as-illustrated structure units, and consequently, enhanced the
resonance of the nanocomposites.
2.3.2 Effect of morphologies of CNFs
Figure 2.5 Influences of microstructures of CNFs on negative permittivity of PEI/CNFs
nanocomposites. The two types of CNFs applied in our studied are illustrated: Cup-stacked CNFs
(CS-CNFs) have circular cross section with large hollow core, whereas herringbone CNFs (HB-
CNFs) have rectangular, hexagonal and heptagonal cross section with solid core. Although the
resonance phenomenon was also found in HB-CNFs filled PEI nanocomposites, the negative
permittivity was restrained, and at most frequencies, the permittivity was positive. Only at the
frequencies around resonance, was the permittivity slightly “dragged” into negative region.
Carbon nanofibers have different microstructures depending on various catalysts and growth
. These different microstructures, like the dimensions and dispersion of CNFs, also
have differing impact on the physical properties of polymer nanocomposites
. For a comparative
study, herringbone CNFs (HB-CNFs), were applied to find out the influences of microstructures of
CNFs on the dielectric constant. Their structure differences have been studied in detail by Yoong-
Ahm Kim et al and Takashi Yanagisawa et al [23-24]
. As shown in Fig. 2.5, the microstructures of
CNFs have a significant effect on the realization of negative permittivity in the PEI nanocomposites.
Although similar resonance phenomena are observed, HB-CNFs seem not to be as effective as CS-
CNFs in leading to negative permittivity. This is probably related to greater number of active edges
in CS-CNFs. Their influences on the electric properties have been found recently by Kyle Ritter et
. Owing to the circular cup-stacked structure with hollow core, the CS-CNFs have much more
active edges (inner and outer) which are highly polarizable 
and consequently enhance the
interfacial polarization between CS-CNFs and polymer chains. At the same time, we do not exclude
the possibility that the non-identical length distribution of the two kinds of CNFs could have some
influences on the negative permittivity, even though both are as-received long CNFs.
2.3.3 Effect of chemical structures of polymers
The interfacial interactions, especially interfacial polarization between polymer and nanofillers,
are significant to dielectric properties of polymer nanocomposites, thus, polyimide (PI) was also
included as a second polymer to study any variations and understand the contributions of a polymer
matrix (chemical structures) to negative permittivity. The results show that the chemical structures
have profound effects on the negative permittivity of polymer nanocomposites. Compared with
ether groups on PEI chains, imide groups strongly promote the negative permittivity in PEI
nanocomposites, probably by interfacial polarization between imide chains and the active edges of
CS-CNFs. The much higher polarity of imide groups, compared with ether groups, tends to generate
higher interfacial polarizability between imide groups and CNFs, and therefore larger amplitude of
negative permittivity. That is why PI nanocomposites with equivalent CS-CNFs loading have much
more distinct negative permittivity (Fig. 2.6). The “lower” negative permittivity in the PEI
nanocomposites should be induced by dilute effect from the ether groups and higher free volume of
phenylene ether groups that usually account for the lower positive permittivity of PEI compared
with that of PI 
. This meaningful result suggests that, when constructing metamaterials with
polymer host/substrate, the polymer component should not only be considered as an insulating
medium, as in most research [11,12,16,17, 27]
, but its chemical structure can also exert great influences
on negative permittivity, and possibly negative permeability. The fact that negative permittivity
appears in organic polymer/non-metallic CNFs composites and epoxy/nanopolyaniline hybrid 
also reveals that metallic element (in metals and ceramics) is not indispensable to the single
negative permittivity materials.
Figure 2.6 Influences of chemical structures of polymers on negative permittivity of
nanocomposites. With the same concentration of CNFs (2.0wt%), the negative permittivity of PI
nanocomposites is a dozen times higher than that of PEI nanocomposites, even close to
PEI/5.0wt%CNF nanocomposites. It indicates the contribution of imide groups to negative
permittivity, while the larger and non-polar phenylene ether groups could depress this phenomenon.
The key structure units of PEI and PI are given aside.
The polymeric “single negative permittivity metamaterial” was successfully fabricated from
organic PEI/PI and non-metallic carbon nanofibers. The determinant structure in polymer
nanocomposites is the 3D network composed of imide chains wrapped on long CNFs (Fig. 2.2),
which is equivalent to the thin metallic wires in J.B.Pendry’s studies 
. Therefore the
nanostructures of the CNFs are essential to the efficient realization of negative permittivity for the
nanocomposites. On the other hand, strong composition dependence effects including chemical
structure dependence of negative permittivity indicate that the magnitude of negative permittivity in
the nanocomposites can be adjusted simply by changing the concentrations of CS-CNFs and
altering the chemical compositions of the polymer chains, such as introducing non-polar groups
with larger free volume.
The discovery of single negative permittivity metamaterials in unstructured polymer
nanocomposites will lead to revolutionary applications due to the great potential in creating
lightweight, conformable polymeric metamaterials. The applications may include novel superlens,
wave guides and band gap materials, and especially interesting, truly drapable invisible cloaks that
cannot be realized by rigid inorganic metameterials as currently foreseen. Particularly relevant
in our study, high performance PEI is popular in microelectronic, aerospace, and semiconductor
industries due to its superior thermal stability, mechanical and dielectric properties, as well as good
solubility and processability. Nano carbon fillers with advanced mechanical and physical properties
have been well recognized as highly efficient modifier for polymers. Given the superior mechanical,
electric and thermal properties of PEI/PI nanocomposites filled with a nano carbon filler, we can
envision a blooming vista of new generation opto-electronic devices, manifold structures and the
transport of energy, people and information, heretofore unimagined. It is also believed that such
foundational knowledge can lead to greater understanding of the prerequisites for achieving organic
metamaterial films with negative permittivity and negative permeability through demonstrating
model equivalence between natural constructions and artificial assemblages.
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The Strong Influence of Carbon Nanofiber Network Variability on the
Pronounced AC Conductivity of the Polyetherimide Composite Films
In today’s world of spiraling increases in demands for electrical power, AC power is receiving
greater interest for its flexibility, ease of transmission and efficiency. In some cases, as in aerospace,
power from DC sources is converted to AC for just those benefits. Thus, as nanocomposites are
increasingly sought after for electrical as well as other functionalities, it is necessary to gain a
comprehensive knowledge of the properties of carbon nanofillers with different dispersion states in
a polymer matrix – exposed to an AC electric field - for further development of conductive polymer
composites. In this study, we report the very pronounced and distinctive post-percolation AC
conduction phenomenon of polyetherimide nanocomposite films containing CNF agglomerates.
Results indicated that configuration variability of carbon nanofiber (CNF) networks had dramatic
effects on the AC conductivity of polyetherimide composites.
Adding conductive nanofillers, such as metal nanoparticles 
, conducting polymer
, carbon nanofillers [4-6]
, etc. to a polymer matrix is an effective and promising
means to obtaining conductive polymer composites for various practical applications such as
sensors, electrostatic charge dissipation, field emitter and electromagnetic interference shielding, etc.
In particular, fibril carbon nanofillers, such as carbon nanofibers and carbon nanotubes (CNTs) have
gained heightened interest from both industry and academia, owing to their versatile physical
properties and superior mechanical performance. Such a polymer-carbon nanofiller composite
system usually shows insulator-conductor transition at a critical loading lever, that is, so-called
percolation conduction [7-8]
, which is mainly dominated by dispersion and distribution of these
conductive nanofillers in polymer matrix besides their intrinsic electronic properties.
Because of greatly increased specific surface area and surface energy in the nanoscale, the
aggregation of nanoparticles quite naturally occurs. For CNFs and CNTs, the entangled fibrous
network also results in the formation of agglomerates. However, in past few decades, although
tremendous efforts have been directed towards the relationships between dispersion state of
nanofillers and percolative conduction of conductive polymer composite [7-10]
, the knowledge of the
influences of conductive nanofiller agglomeration states on electrical conduction of polymer
composites is still very limited.
On one hand, most studies reported have been numerical analyses [10-13]
, and for the convenience
of calculation and stimulation, they often do not reflect the real structural information of the
conductive polymer composites 
, such as the geometry of agglomerates, structures and properties
of polymer matrices, or even the geometry of the test samples. Therefore, the simulated results
seemed inconsistent with experimental findings at times. For example, Pegal 
et al. found a
certain CNT agglomeration could enhance the formation of percolation network when studying
polycarbonate/multi-walled carbon nanotubes composites, while Hu 
et al. concluded that
aggregates of CNTs can increase the percolation value based on 3-D resistor network modeling.
On the other hand, how the agglomerates in polymer matrix behave in an AC electrical field has
rarely been investigated 
, experimentally or theoretically. Presently, AC power driven facilities
are getting increasingly important in practical applications. Thus, the study of the electrical
response of carbon nanofillers with different dispersion states within a polymer matrix exposed to
an AC electric field is profoundly critical to the development of future conductive polymer
composites. In this study, we will report the very pronounced and distinctively responsive post-
percolation AC conduction phenomenon of our polymer nanocomposite films containing CNF
agglomerates, and its susceptibility to network configuration; behavior not previously seen in other
conductive polymer nanocomposites.
Fig. 3.1 TEM image of cup-stacked carbon nanofiber
This study is based on a series of polyetherimide (PEI)/CNF composite films with ca. 50µm
thickness, prepared by a solution processing method in the presence of dicholoromethane.
Polyetherimide (ULTEM 1000) was supplied by Sabic Innovative Plastic Inc. Cup-stacked carbon
nanofibers (Pyrograf® III, with conductivity =1.8×10-3
S/cm, Fig. 3.1), diameter range of 60nm-
150nm and length range of 30um-100um, were purchased from Applied Science Inc. The detailed
film preparation procedures were introduced in Ref. 
3.2.2 Control of configuration of CNF network
Figure 3.2 Optical microscope photos and SEM images of PEI nanocomposite films: Composite A.
PEI/3.0wt% CNF (untreated) composite; Composite B PEI/3.0wt% CNF (probe- ultrasonication
treated); Composite C. PEI/3wt% CNF (Bath- ultrasonication treated)
In order to obtain distinctly different dispersion states of CNFs in PEI matrix, three pre-
treatment approaches were applied (Fig. 3.2): Composite A: no pre-dispersion, in which the CNFs
mainly exist in the form of big agglomerates composed of strongly entangled long CNFs, with small
amounts of individually dispersed CNFs; Composite B: 1hr probe-ultrasonication (Branson 450)
pre-treatment, in which the probe-ultrasonication has 80 watt energy output, and is strong enough to
break the long CNFs (therefore, the dispersion of CNFs in Composite B was greatly improved, and
most CNFs were uniformly dispersed in PEI with few loose and small agglomerates existing, and
consequently, the film samples of composite B exhibited a quasi-homogeneous black color as
shown in Fig. 3.2); Composite C: 9 hour bath ultrasonication (Branson 1510) for CNFs. With the
pre-treatment of low intense bath sonication, the CNF agglomerates were gradually disentangled
and formed a greater number of smaller ones without reducing the lengths of the CNFs, which
endowed Composite C with microstructures in Fig. 3.2 C.
3.2.3 Microstructure analysis
The size distribution of CNF agglomerates is summarized in Table 1 by analyzing optical
microscopy images (OM, Olympus BX51) and scanning electron microscopy (SEM, FEI F200)
images. Since all the CNFs were well embedded in the PEI matrix, in order to investigate the
morphologies of CNF agglomerates and dispersion state of CNFs by SEM, the film surface was
rinsed with dichloromethane to expose the bare CNF surfaces by removing the PEI film surface,
before observation. In Composite B, the CNFs were well dispersed in the PEI matrix, and there
were almost no agglomerates observed by either OM or SEM that were larger than 10µm (mean
diameter of an agglomerate with irregular shape). It is assumed that all CNF agglomerates
(including single CNFs) were smaller than 10µm. To verify this presumption and guarantee the
accuracy of our subsequent analyses, sufficient samples were studied.
3.2.4 Electrical properties analysis
DC and AC conductivities as well as dielectric properties of PEI/CNF composite films were
investigated. DC conductivity was measured by Keithley 6517A Electrometer equipped with
Keithley 8009 test fixture. AC conductivity and dielectric properties measurement were conducted
on a Universal Dielectric Spectrometer BDS 20 from Novocontrol Inc. For AC conductivity and
complex permittivity, only their real part was discussed in this paper.
3.3 Results and discussion
3.3.1 Electrical properties under DC and AC electrical fields
Previous studies 
on the dielectric properties of PEI/CNF composite films have already
revealed the contribution of CNF agglomerates to the unusual negative permittivity in organic
polymer composites (Composite A, Fig. 3.3A). In such a representation the CNF agglomerates
were treated as a continuous 3-D network, and the negative permittivity is interpreted by utilizing
J.B.Pendry’s 3-D metallic thin wires model 
. The continuity of network structures has been
shown to be destroyed by high intensity probe ultrasonication, thus the negative permittivity, as
well as the oscillation of frequency-dependent permittivity, under that severe treatment, disappeared
(Composite B) 
Furthermore, we found unusually high AC conductivity from Composite A films, as shown in
Fig. 3.3B and Fig. 3.4, which is suspected to be related to the negative permittivity. This is
consistent with current thinking that negative permittivity is a consequence of high conductivity.
We observed that the higher AC conductivity of Composite A (ca. 10-4
S/cm) as compared to
that of Composite B (ca.10-7
S/cm), occurred with the same CNF loadings (3.0wt% in Fig. 3.3).
Since all composites studied showed constant AC conductivity across the whole frequency range,
for our comparison between AC conductivity and DC conductivity, it is only necessary to evaluate
Figure 3.3 (A) Real part of permittivity of PEI/ 3.0wt% CNF composite films; (B) Comparison of
DC conductivity and AC conductivity at 10-2
Hz of PEI/CNF 3.0wt%CNF composite films
Composite A with 3.0wt% CNFs
Composite B with 3.0wt% CNFs
Composite B with 10.0wt% CNFs
Composite C with 3.0wt% CNFs
PEI Composite A Composite B Composite C
AC conductivity at 10-2 Hz
AC conductivity at a single frequency, which we chose as 10-2
Hz (Fig. 3.3B). Typically, the AC
conductivity at such a low frequency is equal to or close to DC conductivity, which is the reason
that many researchers [7,15]
use low frequency AC conductivity to analogously study DC conduction
0 1 2 3 4 5 6 7 8 9 10 1110
Concentration of CNFs (wt%)
DC conductivity of Composite A
AC conductivity at 10-2Hz of Composite A
DC conductivity of Composite B
AC conductivity at 10-2Hz of Composite B
Figure 3.4 DC and AC conductivities at 10-2
Hz of PEI/CNF composite films with different
According to the above results, DC conductivity of Composite A was expected to be higher than
that of Composite B, in concert with the AC conductivity. However, surprisingly, in our study, the
DC conductivity of Composite A was found to be not higher, but even slightly lower than that of
Composite B (Fig. 3.4). Thus, the well-recognized relationship between DC conductivity and low
frequency AC conductivity, being equal or close, may not be appropriate to Composite A filled with
CNF agglomerates. After comparing the DC conductivity (2.78×10-17
S/cm) and low frequency AC
S/cm) of pure PEI films, we assumed that the differences of the two types
of conductivities of within two orders is normal in our study. Thus, we can find that, in the
Composite B with well dispersed shorter CNFs (Fig. 3.3B and Fig. 3.4), the AC conductivity at10-
2Hz is only dozens of times higher than the DC conductivity, while for Composite A shown in Fig.
3.3B, the difference is as high as 4~5 orders, which has not been previously reported in any polymer
3.3.2 Microstructures and correlations with electrical properties
Figure 3.5 Statistics of size distribution of composite A and composite C
0 20 40 60 80 100 120 1400
Mean Diameter (m)
20 40 60 80 100 120 1400
A comprehensive study of CNF agglomerates was conducted by investigating dielectric and
electrical properties of Composite C with the same CNF loading. The 9-hour low intense bath
sonication could only decrease the size of CNF agglomerates, but not change the length of the CNFs,
which could help us understand the influences of the size of CNF agglomerates on the electrical and
dielectric properties of PEI/CNF composite films. Fig. 3.5 gives the size distribution (mean
diameter) of CNFs agglomerates of composite A and Composite C, which was quantitatively
summarized in Table 3.1. In Composite C, over 90% of CNF agglomerates are smaller than 10µm
compared with ca. 60% in Composite A. Meanwhile, the CNF agglomerates bigger than 20µm
dramatically decreased. The dielectric properties of Composite C films did not show negative
permittivity in the whole frequency range (Fig. 3.3A), which is the same as Composite B.
Correspondingly, the AC conductivity at 10-2
Hz of Composite C films is just a few dozen times
higher than the DC conductivity. This result may reveal that only CNF agglomerates with some
large size levels (such as bigger than ca. 20µm in our composite films) account for the negative
permittivity and high AC conductivity of the nanocomposite films. The disappearance of negative
permittivity can be explained using J. B. Pendry’s metallic thin wires model 
: the breakage of
CNF agglomerates to smaller ones is also a loss of continuity of the CNF 3-D networks.
Furthermore, even though there are still a few CNF agglomerates larger than 20µm in Composite C,
they are not identical to those in Composite A, owing to the loosened entanglement structure caused
by bath ultrasonication. Thus, this composite may not have the same contribution to the special
dielectric and AC conductivity of the PEI/CNF composite films as those in Composite A without
Table 3.1 Size distribution of CNF agglomerates in different composites
In Fig. 3.4, the concentration dependence of both DC conductivity and AC conductivity at 10-2
Hz of Composite A and Composite B are given. On one hand, in Composite A, the remarkable
difference in DC conductivity and AC conductivity at 10-2
Hz does not show obvious concentration
dependence. For various compositions, such a difference is always around 4~5 orders. This may
indicate that such a phenomenon is mainly derived from the 3-D structures of CNF agglomerates,
and is irrelevant to the number of CNF agglomerates. This is inconsistent with the concentration
dependence of negative permittivity reported on our earlier work 
: the absolute value of negative
permittivity increases with CNF concentration. On the other hand, in Composite B, we increased the
loading levels of probe ultrasonication treated CNFs up to 10 wt % CNFs. It is assumed that much
more conductive paths can be formed in PEI matrix at such a high loading level, and consequently
can result in higher electrical conductivities. However, the results showed that the difference
between DC and AC conductivity in this Composite A is still limited to two orders, which further
reveals the contributions of strongly entangled network structures of CNF agglomerates to the
dielectric and electrical properties of Composite A films.
3.3.3 Mechanism analysis
We assumed that the higher AC conductivity of Composite A is associated with negative
permittivity as well as the oscillation of permittivity, and accordingly is associated with the 3-D
< 10µm >20µm >30µm
Composite A 60.6% 17.8% 9.5%
Composite B ~100% - -
Composite C 90.5% 3.5% 1.9%
networks of CNF agglomerates, which can be further understood by classic Drude model of electron
conduction as well as the dielectric nature of AC conductivity.
The Drude Model of electrical conduction is depicted by Equeations (3.1), (3.2) and (3.3).
J E (3.1)
where J is current density; E is electric field; e is charge of an electron; n is number of conduction
electrons; μd is drift mobility of electrons, defined as:
where me is electron mass; τ is relaxation time that is related to the mean time among electron
The Drude model reveals that the electrical conductivity directly depends on the number and the
drift mobility of electrons. The higher conductivity in AC electrical field may suggested that the
drift mobility of electrons in the AC electrical fields increases, because the AC conductivity of
composites A did not show obvious dependence on the concentration of CNFs which corresponds
to the concentrations of electrons. The increased drift mobility can be further understood by the
dielectric nature of AC conductivity.
For dielectric materials, AC conductivity is a kind of dielectric property related conductivity,
and has a complex number form, as shown in Equation (3.4).
0* ' " 2 ( * 1)i i f (3.4)
In general, its real part, σ’(f), relates to dielectric loss by Equation (3.5),
0' 2 "f (3.5)
Frequency / Hz
Figure 3.6 Dielectric loss of PEI / 2.0 wt% GNP nanocompsites with different CNF network
From Equation 3.5, the high AC conductivity can be seen a result of high dielectric loss. Fig. 3.6
shows the dielectric loss of all three composites studied in this work with a CNF concentration of
2.0wt%, all of which have similar DC conductivity of ~ 10-9
S/cm. Obviously, the composite B with
completely uniformly dispersed CNFs has a dielectric loss as low as pure PEI, while the dielectric
loss in composite C with smaller and loosely entangled CNFs is slightly higher than pure PEI.
Contrarily, Composite A exhibits an extremely high dielectric loss in the entire frequency range,
furthermore, this high dielectric loss has distinct frequency dependence: the dielectric loss goes up
with the decrease of frequency. This frequency dependence did not appear in PEI and Composite B,
and Composite C exhibits slight frequency dependence at low frequency range, possible suggesting
the relationship between the frequency dependent dielectric loss and CNF agglomerates, especially
the strongly entangled CNF agglomerates.
For polymer nanocomposites filled with carbon nanofiller, the interfacial polarization is
determinant to the dielectric properties of the composites, which is strongly related to the motion of
electrons in the interphase 
. In these composites, the dipoles are mainly formed by the displaced
centers of positive and negative charges as a result of short range motions of electrons. The high
dielectric loss is usually caused by intense polarization of dipoles in dielectric materials, in other
words, the high mobility of electrons could enhance the polarization, and consequently lead to high
dielectric loss. Therefore, the high dielectric loss could suggest high electron mobility in this case.
For Composite A films, the special polarization of strongly entangled CNF agglomerates in PEI
matrix accounts for the oscillation of the absolute values of permittivity around zero (Fig. 3.2A) as
well as extremely high dielectric loss, which reflects more energetic electron motions than the other
two composites in AC field, and consequently results in higher AC conductivity. Furthermore, this
kind of special polarization process and energetic electron motions in Composite A films should be
related to the continuous 3-D network structures of CNF agglomerates, which endow PEI/CNF
composite films with negative permittivity, analogous to the continuous 3-D network of metallic
thin wires 
. According to the authors’ knowledge, similar phenomenon has not been reported in
any polymer nanocomposites showing aggregation of carbon nanofillers, and will require an in-
depth study to define the mechanisms involved.
In conclusion, we studied the post-percolation DC and AC conductivities of PEI/CNF composite
films with distinctly different dispersion states. The greatly increased AC conductivity (represented
by the conductivity at 10-2
Hz) of the PEI/CNF composites containing relatively large and strongly
entangled CNF agglomerates may possess new electrical conduction behaviors/mechanisms, and
may reveal the distinctive response of electrons in such strongly entangled CNF networks in
polymer films to an AC field, as compared to a DC field, and which also can result in negative
dielectric response. This requires further and more systemic studies. This finding will stimulate and
fertilize studies to further current knowledge on fabrication and properties of conductive polymer
composites. It holds the promise of promoting development of new conductive polymeric
films/membranes with novel and beneficial properties that will widen the applications of AC
conduction to supplant those power systems still relying on, and subject to the limitations of, DC
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Effectual Dispersion of Carbon Nanofibers in Polyetherimide Composites and
Their Mechanical and Tribological Properties
The use of proliferation of nanotechnology in commercial applications is driving requirements
for minimal chemical processing and simple processes in industry. Carbon nanofiber (CNF)
products possess very high purity levels without the need of purification processing before use and
are in growing demand for this quality. Polyetherimide (PEI) has excellent mechanical and thermal
performance, but its high viscosity makes its nanocomposites processing very challenging. In this
study a facile melt mixing method was used to fabricate PEI nanocomposites with as received and
physically treated CNFs. The dispersion of CNFs was characterized by scanning electron
microscopy, transmitted optical microscopy and electrometer with large area electrodes. The results
showed that the facile and powerful melt mixing method is effective in homogeneously dispersing
CNFs in the PEI matrix. The flexural and tribological characteristics were investigated and the
formation of spatial networks of CNFs and weak interfacial bonding were considered as competitive
factors to enhanced flexural properties. The composites with 1.0 wt% CNFs showed flexural
strength and toughness increased by over 50% and 500%, respectively, but showed very high wear
rate comparable to that of pure PEI. The length of the CNFs also exerted great influences on both
mechanical and tribological behaviors.
Polyetherimide (PEI) is a type of high performance thermoplastic with high modulus and
strength, superior high temperature stability, as well as electrical (insulating) and dielectric
properties (very low dielectric constant) [1,2]. It performs successfully in aerospace, electronic and
other applications under extreme conditions. However, the pure PEI has some disadvantages, such
as its brittleness, and wear rate which limits its applications. Thus, an appropriate modification of
PEI that can widen its applications would be welcomed by many industry sectors.
Although many applications of polymer materials involve situations where the tribological
properties are significant to the performance and service life of polymer materials, with the
exception of ultrahigh molecular weight polyethylene (UHMWPE) and polytetrafluoroethylene
(PTFE), most polymers do not possess satisfactory tribological properties. Generally, great amounts
of glass fibers (GF) [3, 4], carbon fibers [5,6] and carbon fabrics [7,8] must be added to the polymer
matrix to improve their wear resistance. For PEI, which is becoming increasingly important in
industrial applications such as airplane interiors, it is necessary to study the tribological properties
of PEI based composite materials. J. Bijwe [1,6-8], A.P. Harsha  and J. Viña  et al. did
comprehensive research on the tribological properties of glass fibers, carbon fibers and their fabrics
reinforced PEI composites under various wear conditions, such as adhesive wear, abrasive wear,
erosive wear etc. Generally, higher loading of these microfillers could lead to better wear/friction
resistance [3, 6, 8]. In many studies, the loading level was up to ca.80vol%. Besides the
concentration of the fibers and fabrics, their orientation also affects the wear and tribological
properties [5, 6]. In PEI/80vol% long GF composites, the tribological properties were best when the
GFs were parallel to sliding direction in adhesive wear mode, while worst when the GFs were
perpendicular to the sliding direction. A.P Harsha studied the erosive properties of PEI/CF
composites and found the erosive rate was highest at 60O contact angle. For carbon fabric
reinforced PEI composites, the fabric in the direction normal to the sliding surface resulted in higher
friction coefficient of the composites. Moreover, the tribological properties also strongly depend on
the wear modes. J.Bijwe et al  studied the wear properties of PEI/GF composites in four different
wear modes. The results showed the composites had extremely good wear resistance in fretting and
adhesive wear modes, but very poor in abrasive and erosive wear modes. Other factors, such as
counterface materials, temperature and sliding speed also exert significant influences on the
tribological properties. In all these PEI composites in tribological applications, a great amount of
microfillers was needed, which can result in the loss of some unique properties of PEI including low
mass density, low dielectric constant, and toughness, and also increase the difficulty in composite
processing. These problems have been shown to be alleviated by adding very small amounts of
nanofillers - rather than microfillers - to the PEI matrix.
Carbon nanofibers (CNFs) [9,10], which in the past few years have been extensively applied in
thermoplastic and thermoset composites, are capable of enhancing many polymeric materials due to
their attractive performance and cost characteristics. CNFs not only have very high strength and
modulus along the axial direction, but also exhibit a structural flexibility due to the weak van der
Waal forces between the stacked graphene layers. CNFs have an attractive set of characteristics for
industry in that they are relatively low cost when compared to carbon nanotubes, and commercial
CNF products possess very high purity levels without the need of purification process before use.
They inherently have many free edges so functionlization is readily accomplished without resorting
to damaging treatments as with carbon nanotubes. These characteristics are very attractive for
industry which in general is seeking reductions or eliminations of chemicals involved in
development of their nanocomposite products.. At the same time, CNFs have high electrical and
thermal conductivity leading to their desirability for use in multi-functional polymer
nanocomposites. The CNFs with their graphitic structures also show promising potential in
improving tribological properties of polymer materials by various wear mechanisms such as acting
as lubricant, reinforcing transfer films and breaking up CNFs, etc [10,11]. Thus far, CNFs have
been successfully used to modify the wear properties of ultrahigh molecular weight polyethylene
(UHMWPE), polyphenylene sulfide (PPS) and polytetrafluoroethylene (PTFE), with the loading
level lower than 5wt%, while maintaining other properties. . Therefore, there is great potential for
the CNFs to be used as cost-effective nano-additives for the development of high performance PEI
nanocomposites without chemicals involved in the processing, which is very attractive for industry.
Nanotechnology with the incorporation of CNFs may meet the demand of chemical reduction
programs that are increasingly being promoted in the industrial sector.
PEI has very challenging processability because the extremely high viscosity of PEI also
severely impedes the uniform dispersion of nanoparticles. Even the solution method, a commonly
used mixing method effective for many polymer materials, cannot disperse CNFs or CNTs well in
PEI. In order to solve the dispersion issue of nanofillers in such a high viscosity polymer matrix,
various approaches have been attempted, including ultrasonication treatment , chemical
treatment , and ultrasound-assisted extrusion . All these showed greatly improved
dispersion of nanofillers in the polymer matrix; however some disadvantages cannot be overlooked,
such as the degradation of the intrinsic properties of the nanofillers due to this severe treatment,
high cost and complicated procedures of treatment, and instrumental operation. All of these reasons
make PEI nanocomposites difficult to fabricate in volume. To be applicable to commercial
applications the materials must have a minimal set of processes and those must be robust and
insensitive to minor parametric perturbations.
In this study, commercial graphitic carbon nanofibers, Pyrograf III ®, in the pristine form and
physically treated format, were used to make PEI nanocomposites. The fiber lengths were cut
shorter by the physical treatment through a high power sonication approach. The dispersion and
mechanical and tribological properties of the PEI/CNF composites were investigated. The even
dispersion of CNFs in PEI matrix was realized by direct melt mixing on a twin screw extruder at an
elevated temperature (3700C). The flexural testing, dynamic mechanical analysis testing and pin-on-
disk sliding wear testing were applied to evaluate the performances of the composites. The
influences of concentration and fiber length of CNFs were studied. Furthermore, the possible
relationship between flexural properties and tribological properties are considered.
4.2.1 Preparation of polyetherimide/carbon nanofiber composites
Polyetherimide (ULTEM 1000) powder was supplied by SABIC Innovative Plastic Inc. CNFs,
Pyrograf III ®PR-HHT-24, with diameter of 60 -150nm and length of 30µm-100µm was provided
from Applied Science Inc. After drying, CNFs were pre-mixed with PEI with different mass ratios
as shown in Table 4.1. The mixture of PEI and CNFs was extruded on a twin screw extruder
(Leistritz) at a temperature of 370 ºC to obtain PEI/CNF nanocomposites. The extrudates were
compression molded (Carver Model 3693) to 3mm thick panel samples at 315 ºC for flexural testing
and DMA testing.
Here, two types of CNFs were used to study the influences of lengths of CNFs on properties of
PEI/CNF composites: pristine CNFs without any treatment and ultrasonically treated CNF. The
ultrasonic treatment of CNF was conducted on a W-385 Heat Systems ultrasonicator (Branson.co.)
for 1 hour in the presence of acetone as dispersant. The power level was 70w. After treatment, the
solution was placed under an evaporation hood to allow the acetone solvent to evaporate leaving the
treated CNFs to be collected for use. Our previous work has already showed that the ultrasonically
treated CNF have shorter length .
Table 4.1 Composition of PEI/CNF composites
PEI Pure PEI
C-0.5L PEI + 0.5wt% CNF (pristine)
C-1.0L PEI + 1.0wt% CNF( pristine)
C-3.0L PEI + 3.0wt% CNF(pristine)
C-1.0S PEI + 1.0wt% CNF (ultrasonically treated)
4.2.2 Three point flexural testing
An Instron 4466 2 kip electromechanical universal test machine was used for the 3-point
flexural test for the nanocomposites. The loading rate was 2.032mm/min. The specimen dimensions
were 50mm (length) x 10mm (width) x 3.0mm (thickness), and the span between two testing
supports was 38.1 mm. For each composite, at least five samples were tested.
4.2.3 DMA measurement
The dynamic mechanical analysis was conducted on a Tritec 2000 DMA test machine in single
cantilever mode. The displacement of each specimen was 0.01 mm with a frequency of 1 Hz. The
heating rate was 5 °C/min ranging from 30 °C to 300 °C. The specimen dimensions were 50mm
(length) x 10mm (width) x 3.0mm (thickness)
4.2.4 Sliding wearing test
Specific volumetric wear rates were determined using a custom-built pin-on-disk apparatus with
a vertical 1020 carbon steel disk under dry testing conditions, with details shown in Section 4.3.4.
Wear testing samples had dimensions of about 10mm x 10mm x 3mm cut from compression
molded panels. Wear testing was performed at 180 rpm at room temperature, with an effective disk
radius of 6.5cm
Specific wear rates were calculated by the mass loss of the sample being tested, using the
following equation (4.1):
where w is the wear rate, ∆m is the change in mass, ρ is the density of the material, F is the normal
force, and d is the linear sliding distance. More details of wear test can be found in ref. 
4.2.5 Microstructure analysis
Scanning electron microscopy (FEI 200F) was applied to study the morphology of the fracture
surfaces and wear surfaces of PEI/CNF composites with an accelerating voltage of 30KV. A thin
gold layer was coated on the sample surface before SEM observation. We also used the transmitted
light microscopy (TLM) method to study the dispersion of CNFs by grey scale analysis - a method
developed by Hattum et al . Furthermore, according to the direct relationships between
electrical conductivity and dispersion state of fibrillar conductivity fillers in the matrix and the DC
conductivity behavior of the PEI/CNF composites were investigated by Keithley 6517A with
Keithley 8009 test fixtures. The combination of SEM, TLM and conductivity measurement ensures
a full understanding of the dispersion state of CNFs in this study. For both transmitted light
microscopy and DC conductivity resistivity measurement, thin films (ca.50 μm) of the composites
were used. The preparation of film samples by solution method has already been introduced in ref.
4.3 Results and discussion
4.3.1 Dispersion of carbon nanofibers in PEI matrix
Figure 4.1 SEM images of fracture surface of PEI/CNF (as-received) composites
In our study, the CNFs were dispersed in the PEI matrix simply by strong shear force from the
extruder at a high temperature. The dispersion of the CNFs in the PEI matrix was systematically
studied by SEM combined with DC conductivity of thin films of PEI/CNF composites. According
to the SEM images of the fracture surface morphologies in Fig. 4.1, it can be seen that the pristine
CNFs evenly disperse in PEI matrix without any visible agglomerates, even at 3.0wt% CNF loading
- the highest loading level in our study. For the optical microscopy study, thin films with thickness
around 50 ± 5μm were prepared by dissolving the extrudates of PEI/CNF composites in
dichloromethane without any other treatment and then cast coating. The homogeneous grey scale
of the micrographs obtained by transmitted light microscopy further reveals the high dispersion
quality of CNFs by melt mixing  (Fig. 4.2c).
Figure 4.2 DC resistivity of thin films of pure PEI and PEI/CNF composites (a) as well as the
Micrographs of transmitted light microscopy for thin films of PEI/CNF composites (b and c)
In addition to the microscopy study, the homogeneous dispersion of CNFs is also strongly
supported by the evaluation of DC resistivity of the same thin films of PEI/CNF composites as
shown in Fig. 4.2(a). It should be noted that DC conductivity of thin films is an effective approach
to evaluating the overall dispersion state of conductive fillers in polymer matrix at moderate loading
(b) DS-3.0L: PEI+3.0wt% CNF (Bath sonication)
PEI C-0.5L C-1.0L C-3.0L C-1.0S DS-3.0L
by melt mixing
by bath sonication
(c) C-3.0L: PEI+3.0wt% CNF (Melt mixing)
level that has been rarely reported. First of all, in conductive polymer/CNF composites, the
dispersion and distribution of CNFs strongly influence electrical properties of the polymer
composites . The aggregation of CNFs can form a local conductive network which may initiate
percolative conduction of the polymer composites. Because of the microscale thickness of thin films,
the conductive path will much more easily form in the thickness direction when the aggregation
and/or effective contact among conductive filler occurs. That indicates the DC conductivity in the
thickness direction of thin films will be very sensitive to the difference in dispersion state. Secondly,
general microscopy studies only focus on an extremely small point of composite sample, especially
Transmission Electron Microscopy (TEM), in this case, sufficient amount of samples have to be
observed to get overall dispersion information. However, in our study, with the large area electrodes
(with diameter of 50 mm) correspondingly big samples, were applied, thus, just a single
measurement can produce good overall dispersion information and quickly detect the existence of
small agglomerates in the thin films by perceiving significant changes of conductivity/resistivity as
compared with pure polymers. For example, in PEI/3.0wt% CNF (bath sonication) composites (DS-
3.0L), the CNFs were dispersed by low intensity bath sonication for 9 hours. From micrograph of
the thin films in Fig. 4.2(b), CNF agglomerates obviously exist. Correspondingly, the resistivity of
the thin films for this composite show a dramatic reduction compared with pure PEI matrix from
ohm·cm to 108
ohm·cm. Contrarily, at a nominal loading level, the homogenous dispersion of
conductive filler has less effect on the resistivity in the thickness direction in thin films, because of
the failure in forming effective conductive networks. This is the rule we used in this study to
evaluate the dispersion state of CNFs in PEI matrix. The DC resistivity measurement results show
that by melt mixing on the extruder, the addition of CNFs only decreased the DC conductivity of
pure PEI slightly from 1016
ohm·cm to 1015
ohm·cm for the entire set of composites made by melt
processing (Fig. 4.2(a)). The extremely high resistivity of PEI/CNF composite films, close to that of
pure PEI, indicates the very uniform dispersion of CNFs in the composites but a lack of percolation
network among CNFs (Fig. 4.2(c)). Based on our systematic study on dispersion of CNFs, it can be
considered that the direct melt mixing at the temperature of 3700C is sufficient to achieve
homogenous dispersion of CNFs in PEI matrix without any special interface modification or
ultrasonication assisted approaches.
4.3.2 Flexural property analysis
According to ASTM D790, the flexural stress, flexural strain, and flexural modulus are
determined by following equations:
Flexural stress ζf:
ζf = (3PL) / (2bd2) (4.2)
where P is the load applied, L is the support span, b is the width, and d is the thickness of the test
Flexural strain εf :
εf= (6Dd) / (L2) (4.3)
where D is the deflection of the specimen.
Flexural modulus Ef:
Ef = (L3m) / (4bd
where m is the initial slope of the straight-line segment of the load-deflection curve.
Figure 4.3 Flexural stress-strain curves: (a) Concentration dependence of flexural properties; (b)
Influence of CNF length on flexural properties
The flexural stress-strain curves are given in Fig. 4.3. Here, the flexural strength (ζfB) and
flexural strain at break (εfB), as well as flexural modulus, are summarized in Table 4.2. The area
under the stress–strain curve represents the energy consumption before the sample breaks and is
typically applied to evaluate the toughness of materials (energy to fracture), which is also given in
Table 4.2. Because the test sample of composite C-1L did not break during the flexural testing as
shown in Fig. 4.4, the flexural strength was replaced by the yield strength determined from the
flexural stress-strain curve (Italics), and the flexural strain at break was replaced by the maximum
strain the instrument can reach during flexural testing. The real flexural strain at break and the area
under the stress-strain curve for composite C-1L should be higher than the values in Table 4.2.
0.00 0.05 0.10 0.150
0.00 0.05 0.10 0.150
Table 4.2 Flexural properties of PEI/CNF composites
Composite Ef (GPa) ζfB (MPa) εfB (%)
Area under stress-
PEI 3.81 ± 0.70 161.23 ± 15.81 5.36 ± 0. 80 4.06±0.65
C-0.5L 3.68 ± 0.31 88.74 ±17.10 3.29 ± 0. 13 1.39±0.01
C-1L 4.01 ± 0.26 248.01±10.70 > 14.61 ± 0. 42 >26.37±0.68
C-3L 4.15 ± 0.11 144.01± 16.82 4.10 ± 0.41 2.62± 0.53
C-1S 4.03 ± 0.29 170.63 ± 45.10 4.62 ± 0.80 3.76± 0.38
As expected, the addition of CNFs in the PEI matrix increases the flexural modulus of the
PEI/CNF composites, due to the contribution of the high modulus of the CNFs. However, the
influences of CNFs on the flexural strength and flexural strain at break of the composites are rather
complicated. From Fig. 4.3 and Fig. 4.4, it is noteworthy that, unlike the other PEI/CNF composites
which fractured in a brittle manner, Composite C-1L (PEI/1.0wt%CNF (pristine)) deformed
plastically showing a predominant yield point and would be expected to fracture in a ductile manner
if sufficient testing capability were available. Initially it was thought that this behavior was the
result of an artifact of the mechanical testing. However as can be seen in the other areas of
evaluation, the composite C-1L also produced stridently different results in sliding wear. Combined
with the study of Composite C-1S (PEI/1.0wt%CNF (ultrasonically treated)) in Fig. 4.3, it is
assumed that the toughening of PEI/CNF composites is associated with both the concentration and
length of the CNFs.
Figure 4.4 Image of test samples after flexural testing
During flexural testing, the bending moment generated within the test sample induces
simultaneous tensile stress and compressive stress. Generally, failure of test samples in flexural
testing happens when the applied tensile stress/compressive stress is greater than the yield stress of
the material. For PEI, its compressive strength (22000 psi) is higher than tensile strength (16500 psi)
according to the data offered by SABIC Innovative Plastic Inc.. Thus, the failure of PEI should start
from the tensile face and this is exactly the situation for the brittle PEI/CNF composites. Thus, in
the case of PEI composites, increasing tension resistance of materials will improve the flexural
properties of materials, and vice versa. It is easy to understand that, owing to the lack of strong
interfacial bonding, the tensile stress cannot be effectively transferred between PEI and CNFs, and
consequently results in dramatically reduced flexural strength and flexural strain at break with only
0.5wt% pristine CNFs loading, as shown in Table 4.2. With increasing loading of CNFs, the
interface area also increases which results in more structural flaws within the composites and
further impaired flexural properties. However, at some point the increasing CNF loading also leads
PEI C-1S C-0.5L C-1L C-3L
to the spatial entanglement of unagglomerated CNFs in the composite, which could resist fracture
and increase the plastic deformation . This may account for the sudden brittle-ductile transition
of PEI/1.0wt%CNF (pristine) composites (composite C-1L). Specifically, above a certain loading
level, the effective CNF crosslink forms in the composites, and acts as a reinforcing component.
Figure 4.5 Cryo-fractured surface of PEI/1.0wt% CNF (pristine) composite (a) and
PEI/1.0wt%CNF(ultrasonically treated) composite(b).
When 1.0wt% pristine CNFs are added, the entangled reinforcement effect overrides the
negative effect caused by weak interfacial bonding. The tensile stress the composite can resist is
higher than its yield strength; as a result, the composite C-1L deformed plastically and would
ultimately fracture in a ductile manner. According to Table 4.2, compared with pure PEI, the
PEI/1.0wt% CNF (pristine) composite (composite C-1L) shows 53.9% increase in strength, over
172.6% increase in flexural strain, as well as ~549.5% increase in fracture toughness. The similar
reinforcement and toughening phenomenon due to the spatial crosslinking of 1.0wt% CNFs has also
been observed in PTFE/CNF composites . On the other hand, with further CNF loading to
PEI/1.0wt%CNF(ultrasonically treated) PEI/1.0wt%CNF(pristine)
3.0wt%, the negative effect of weak interfacial bonding prevails, owing to greatly increased
interface area. In this case, the composite C-3L fracture mode returned to the brittle manner with
reduced flexural properties compared with pure PEI as shown in Table 4.2, however, still much
better than those of composite C-0.5L, which could be attributed to some minimal CNF crosslinking
effects as mentioned above.
The study of PEI/1.0wt%CNF (ultrasonically treated) composite further reveals the significance
of concentration and crosslinking of CNFs in the toughening of PEI/CNF composites. Fig. 4.5 gives
the SEM images of fracture surfaces of rectangular samples of Composite C-1L (Fig. 4.5(a)) and
Composite C-1S (Fig. 4.5(b)). Both samples were broken in liquid nitrogen. Different from pristine
CNF filled PEI composite (C-1L), there are almost no long fibers coming out of the fracture surface
of PEI/1.0wt% CNF (ultrasonically treated) composite. Because of the shorter length of
ultrasonically treated CNFs from damage during the process, the crosslinking of CNFs in composite
C-1S is not as effective as that in composite filled with long pristine CNFs in inducing the plastic
deformation of PEI matrix. Therefore, the flexural properties of PEI/1.0wt%CNF (ultrasonically
treated) dramatically declines compared with PEI/1.0wt%CNF (pristine), but are still superior to
those of PEI/3.0wt%CNF (pristine). This result indicates that on one hand, long CNFs leads to
sufficient spatial crosslinking to induce plastic deformation of PEI, but on the other hand, 1.0wt% is
the appropriate CNF loading level for PEI composites, at which the CNF crosslinking is sufficient
to induce plastic deformation while overwhelming the negative influence of poor interfacial
4.3.3 Dynamic mechanical properties
Figure 4.6 Storage modulus (a) and loss modulus (b) of PEI/CNF nanocomposites
As shown in Fig. 4.6, the addition of CNFs exerts great influences on dynamic mechanical
properties of PEI/CNF composites. With only 0.5wt% CNFs loading, unexpectedly, the storage
50 100 150 200 250 3000
50 100 150 200 250 3000
modulus of PEI/CNF composite does not show any increase, which is close to the results of flexural
modulus. However, with increasing loading of pristine CNF in PEI matrix, the storage modulus
dramatically rises in a rather wide temperature range. Particularly, for PEI/3.0wt% CNF (pristine)
composite, its storage modulus is almost 70% higher than that of pure PEI. At the same time, it is
noteworthy that, even at very high temperatures over 2000C, the storage modulus of PEI/CNF
composites is still at a rather high level, which may satisfy the requirements of high temperature
applications, The lower storage modulus of composite C-1S is attributed to the shorter length of
ultrasonically treated CNFs. The same influences of fiber length on modulus have also been
reported in microfiber reinforced polymers in previous studies [19, 21].
The loss modulus is slightly increased after adding CNFs into PEI matrix, which is related to the
frictional motion among CNF and PEI, as well as among CNF themselves. A previous work 
showed that the frictional sliding at the nanotube–polymer interfaces can deliver an order of
magnitude (1,000%) increase in loss modulus of the bulk polycarbonate system with only 2%
weight fraction of oxidized SWNT fillers. However, in our study, the loss modulus does not show
any further significant increase at higher CNF loading, probably owing to the weak interfacial
bonding and resulting less effective frictional sliding among pristine CNFs and PEI. As for Tg , in
our study, a high heating rate (50C /min) was applied during DMA testing, thus, the Tg of pure PEI
is much higher than 2160C supplied by manufacturer. A 25
0C reduction of Tg is observed when the
CNF loading is higher than 1.0wt%. This may be explained by the formation of voids within fiber
reinforced polymer composites. It is well known that Tg is a reflection of mobility of polymer chain
segments. The movement of polymer chain segments needs both sufficiently high temperature
(kinetic factor) and sufficient free volume (space). The formation of significant voids is common in
CNF filled polymer composites , which can also be found in SEM images in Fig. 4.1 and Fig.
4.5. The existence of these voids in the composite supplies more free volume for rigid PEI chain
segment to motion, and results in a decreased Tg
4.3.4 Sliding wear properties
Generally, carbon nanofillers with graphitic structures can act as a solid lubricant. CNFs also
have the ability to contribute increased load-bearing capacity to polymer composites and reinforce
transfer films produced during wear testing [10, 11, 24]. Thus, the addition of CNFs to a polymer
can obtain polymer composites with improved tribological properties showing decreased wear rate.
Figure 4.7 Wear rate of PEI/CNF composites: (a). with respect to sliding distance; (b). at steady
However, our results did not show such a tendency. According to Fig. 4.7(a), the existence of
CNFs does not necessarily decrease the wear rate of PEI/CNF composites. Although the 0.5wt%
and 3.0wt% CNF filled PEI composites exhibit 56.25% and 65.38% reduction in wear rate at steady
state wear (Fig. 4.7(b)), both PEI/1.0wt%CNF composites filled with treated and untreated CNFs
have a wear rate as high as pure PEI. Thus, besides the effects of CNFs on tribological properties
PEI C-0.5L C-1L C-3L C-1S0.001
400 600 800 1000 1200 1400
Figure 4.8 SEM images of wear debris of PEI/CNF composites after pin-on-disk testing
introduced above, CNFs may also have other dominating influences which will be discussed later.
The reasons for the wear properties of PEI/CNF composites can be firstly understood through SEM
images of wear debris. It can be seen that for pure PEI (Fig. 4.8(a)), its debris mainly exists in the
form of large ribbon structures indicating a severe abrasive wear mechanism which usually results
in poor tribological properties of polymer composites. Similar ribbon structures are also observed in
both 1wt% CNF filled PEI composites (Fig. 4.8(c) and 4.8(e)). While in Composite 0.5L (Fig.
4.8(b)) and Composite 3L (Fig. 4.8(d)) CNF filled PEI composites, the wear debris size becomes
much smaller, and the ribbons almost disappear suggesting less abrasive wear occurred in these two
composites and consequently decreased wear rate.
Interestingly, after comparing the flexural properties and wear properties of PEI/CNF
composites, it is found that the wear rate of PEI/CNF composites shows approximately similar
composition dependence as their toughness, that is, the composites with high toughness have lower
wear resistance, which has not previously been found in other polymer composite systems and may
be the predominant reason for the unexpected poor wear properties of 1.0wt% CNF filled PEI
composites in our study. At present, the relationship between mechanical properties (modulus,
strength and toughness) and tribological properties of polymer composites has not been
comprehensively studied and are still unclear. This usually varies with different composite systems
and diverse wear modes. For example, J. Bijwe et al found the mechanical properties and abrasive
wear performance of fabric-reinforced PEI could not be correlated , J. Karger-Koscis et al.
studied the CNF reinforced Santoprene Thermoplastic Elastomer and no direct relationships were
found between tensile properties and tribological properties under different wear testing, while M.C.
Galetz et al.  considered the improvement of wear properties of CNF/Ultrahigh molecular
weight polyethylene was contributed by the increase in matrix modulus, yield strength and hardness
due to the CNFs.
In contrast, in our study, both flexural modulus and storage modulus increases with CNF
loading, while the flexural strength varied conversely to wear resistance (Table 4.2). However, an
approximate relationship between toughness (εfB and area under stress-strain curve) and wear rate
could be postulated, which may be dominated by the deformation mechanism involving complete
energy dissipation in the contact area . As discussed in section 4.3.1, the addition of 1.0 wt%
pristine CNFs to PEI matrix may form a spatial crosslink network which can induce/improve the
plastic deformation of PEI matrix. During the sliding wear test, the spinning disk applied shear
stress to the wear surface of PEI/CNF composites, and the effective crosslinked CNFs in tougher
composites could induce plastic deformation under this stress, resulting in more “pulled/scratched-
out” materials and consequently enhanced abrasive wear (Fig. 4.8(c) and 4.8(e)) .This is why the
PEI/1.0wt% CNF composites did not show dramatically decreased wear rate as did the 0.5wt%
CNF filled composites. At the same time, the positive effect of CNFs on wear rate reduction
maintained the wear rate of Composite 1L at the similar level as pure PEI. The slightly higher wear
rate of Composite 1S compared with Composite1L (Fig. 4.8) is assumed due to the short fiber
length of ultrasonically treated CNFs. Many previous studies have revealed that short fibers
(regardless of the fiber types) [28, 29] were less effective than long fibers in reducing wear rate of
polymer composites. Furthermore, for the other two low toughness composites, the lower wear rate
of Composite 3L should be caused by its higher concentration of CNFs as compared with
The studies on PEI/CNF composites reveal that CNFs can be effectively and homogeneously
dispersed in high viscosity PEI matrix by a melt mixing processing method. The mechanical
properties of PEI composites with non-chemically treated CNFs showed composition dependence
due to two competing influences, nanofiber networking and nanofiber-PEI interface flaw density.
High concentrations of pristine CNFs led to increased interfacial flaw density, accordingly, the
reinforcing effects that should be provided by the high loading of CNFs failed to be manifested in
enhancing the mechanical properties.. The wear rate of the PEI/CNF composites correlates with the
toughness. The fiber length is seen to be essential to both flexural and tribological performance of
PEI/CNF composites, and the short CNFs not being as effective as long CNFs. Further studies are
justified to fully understand and model this phenomenon in order to enhance applications of
This study, revealing that uniform dispersion of CNFs in PEI can be obtained in a melt mixing
method, provides great potential for industry to use CNFs directly in developing nanocomposites. It
demonstrated that high mechanical property enhancement can be achieved with commercial CNFs
(untreated) when the loading levels are well controlled. In addition, when chemical treatments are
necessary, they are minimized in energy, time and chemical usage because CNFs inherently have
many free edges, thus enabling nanocomposites with broader multi-functionalities and lower cost
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Effective Static Dissipation of Bi-layer Thermoplastic Nanocomposites
at Low Nanofiber Loadings
It is challenging to obtain effective electrostatic dissipation performance for nanocomposites at
low nanofiller loadings, which requires low resistivity and low relative permittivity. In this study,
we present a bi-layer structured polyetherimide (PEI)/carbon nanofiber (CNF) composite fabricated
by a facile solution casting method, showing dramatically improved static dissipation properties at a
rather low loading level. Compared to the conventional mono-layer nanocomposite with high static
dissipation rate, the bi-layer structured composite exhibited a 91.4wt% reduction in CNF loading,
while it’s volumetric and surface static dissipation rates were ca. 1000 times and 20 times higher,
respectively. The interface region in the bi-layer structure accounts for such a remarkable
improvement. It is believed that this design could significantly benefit the applications requiring
high antistatic capability of polymeric materials.
Generally, the effective static dissipation for electrostatic discharging (ESD) protection requires
a surface conductivity of 1012
. As is well-known, except for properly doped conjugative
polymers (conductive polymers), all other polymers are excellent insulators with extremely high
volume resistivity and surface resistivity. Thus, the design for ESD protection is crucial to the
application of polymeric materials and has attracted great interest in past decades [1-13]
Conventionally, the mainstream approaches to improve the static dissipation of polymeric materials
include: 1) antistatic coating [2-6]
, 2) blending with conductive polymers [7-9]
, and 3) the addition of
conductive fillers etc.[5,10-13]
. A conductive paint made with polyaniline
(PANI)/dodecylbenzenesulfonic acid (DBSA) dispersion and poly(methyl methacrylate) (PMMA)
was prepared and showed extremely high conductivity when PANI/DBSA: PMMA was 1:4 (w:w)
. Through cationic UV curing
, antistatic epoxy/carbon nanotube coating was fabricated with
dramatically improved surface conductivity at a low loading level (0.1wt %). Ruskenstein et al 
synthesized poly(3-methylthiophene) coated polyurethane with high surface conductivity. The
conductive polymer can not only be used as conductive coating, but also can be blended with other
insulating polymer matrices to improve the static dissipation property, such as highly conductive
PANI complex/polypropylene composites 
and polypyrrole/polyethylene, polypyrrole/
polypropylene, polypyrrole/poly(methyl methacrylate) composites 
. However, for these two
methods, the disadvantages are obvious, owing to the weak bonding (different physical and
chemical characteristics between coating and protected parts), the peel-off of conductive coating
layer is very common after short-term service, that is, the conductive coating is not ideal for long
term antistatic purpose. For the conductive polymers, its poor processibility and mechanical
properties (brittleness) limited their applications. Currently, in-situ polymerization is one of the
most common methods [6,14]
to fabricate conductive polymer composites/blends, thus, it is still
challenging to fabricate bulk parts for static dissipation purpose using conductive polymers.
Addition of conductive fillers, especially conductive nanofillers, is a convenient and efficient
way to obtain the polymer composites with desirable static dissipation properties. With proper melt
and solvent processing methods 
, either bulk parts or thin films can be fabricated. In the past few
decades, many studies on various nanofillers and polymer nanocomposites have been conducted [15-
17]. Great potential for polymer nanocomposites have been shown due to the excellent mechanical
and physical performances of nanofillers. Among various nanofillers, carbon nanofillers, such as
carbon nanofiber (CNF), carbon nanotube (CNT) and graphite nanoplatelet, have succeeded in
improving electrical, dielectric, thermal, and mechanical properties of polymer nanocomposites
with low loading of the fillers. At present, the improvement of surface conductivity/static
dissipation of polymer nanocomposites is also mainly focused on the carbon nanofiller modified
polymer nanocomposites. Yuen et al studied the polyimide/multi-walled CNT (MWCNT)
, and found with the addition of 5phr MWCNT, the surface conductivity could
increase to ca. 108(Ω/cm
 et al. added octadecylamine modified MWCNT to
polypropylene and obtained a surface resistivity of ca.109ohm/sq with 2wt% loading. In these two
cases, the mechanical properties have also been improved. In MWCNT modified EVA and PS
composites, Lim 
found with proper control of processing parameters (rotor speed and cooling
rate), the percolation threshold of surface conductivity can be as low as ca.0.75wt% in EVA
composites and 0.5wt% in PS composites, while the conductivity reached around 105Ω/Sq.
These studies discussed the static dissipation performances by investigating either volume
resistivity, or surface resistivity, or both. Although the improvement in both conductivities was
dramatic, most studies neglected another important aspect affecting the static dissipation, viz.
permittivity. The rate of charge dissipation can be evaluated by:
𝜏 = 𝜌𝜀0𝜀𝑟 (5.1)
where, η is the time constant for charge dissipation rate; ρ is the electrical resistivity (both volume
and surface resistivity), and ε0 is the permittivity of free space as a constant, εr is the relative
permittivity of materials. In contrast to conductivity, relative permittivity represents the ability of a
material to store charges. Consequently, for a good static dissipation performance, not only the low
resistivity, but also the low relative permittivity is required, which has been realized in this study.
Since ε0 is a constant, in the following sections, the calculation will use the product of ρ and εr
However, the addition of carbon nanofillers could significantly increase the dielectric properties.
The relative permittivity of a pure polymer is usually located in the range of 2-10, while the
modification by carbon nanofillers can result in the composites with relative permittivity from a few
hundreds to 104~5
, or even higher[18-20]
. Sui et al prepared PP/CNF composites by melt mixing, and
found the enhancement in relative permittivity from ca. 3 of PP to over 600 below 4000 Hz when
the loading of CNF was up to 5wt% 
. Dang et al 
used pristine MWCNT to modify
polyvinylidene fluoride (PVDF), the relative permittivity increased to around 500 at 102 Hz with the
loading of 3wt% MWCNT. He et al 
added 2.5vol% exfoliated graphite nanoplatelets to PVDF
matrix, the relative permittivity of the resulting composites was as high as 107. Thus, the
simultaneously increased conductivity and relative permittivity by adding carbon nanofillers to
polymer matrix are two competing aspects, limiting the polymer/carbon nanofiller composites for
static dissipation application.
It is found that in mono-layer composites, it is difficult to simultaneously realize low surface
resistivity, low volume resistivity and low relative permittivity at a low loading level.
In this paper, we report a bi-layer-structured polymer nanocomposite with very low filler
loading prepared by facile solution processing method, which show dramatically improved surface
conductivity and maintains the relative permittivity close to the polymer matrix. Polyetherimide
(PEI) was explored as the polymer matrix owing to its outstanding mechanical properties and
thermal stability, and desirable electrical and thermal properties make PEI an important material in
both aerospace and microelectronic applications. CNF with high electrical conductivity and low
cost was used as the conductive filler to modify the static dissipation property of PEI matrix. The
solution casted film/thi sheet samples with two dispersion states (Fig.1) were prepared to fabricate
the bi-layer structured composites. These different dispersion states of CNF were realized by bath
sonication treatment (non-uniform dispersion) and probe ultrasonication (uniform dispersion)
treatment of CNFs. The two composites are specified by N-Composite and U-Composite,
5.2.1 Raw materials
Polyetherimide (ULTEM 1000) powder was supplied by SABIC Innovative Plastics, Inc. Cup-
stacked carbon nanofibers (Pyrograf® III, PR-24-HHT), with diameter of 60nm-150nm and length
of 30um-100um, were supplied by Applied Science Inc.
5.2.2 Preparation of polymer nanocomposites
Two types of CNFs were obtained by using two different ultrasonic treatments: High power
ultrasonication and low power ultrasonication. The high power ultrasonic treatment (probe
ultrasonication) of CNFs was processed on a Branson digital ultrasonicator (Model 450) at 20%
amplitude for 1 hour in the presence of the diluent Dichloromethane. After treatment, the PEI power
was directly added to the CNFs solution. The spin mixer was used for over 12 hrs to accelerate the
dissolving of PEI in Dichloromethane and mix PEI with CNFs. At last, the vessel containing
PEI/CNF/Dichloromethane solution was immersed in water bath for low power ultrasonication
(Branson 1510) for 2 hours. The low power ultrasonic treatment of CNFs was conducted in water
bath (Branson 1510) in the presence of Dichloromethane as dispersant for 9hrs. And then, the PEI
powder was added to the CNF solution. The following steps of nanocomposites preparation are
exactly the same as introduced above.
Both mono-layer and bi-layer-structured composites films/sheets with thickness of 50um to
1mm were cast coated on a glass plate. After the solvent completely evaporated, the coated glass
plate was immersed in cold water for a while in order to separate the samples from the glass plate.
5.2.3 Morphology and properties analysis
The microstructures of PEI/CNF (probe ultrasonication) composite and PEI/ CNF (bath
sonication) composites were analyzed by Optical Light Microscope (Olympus BX51) at the
magnitude of ×50. The difference in dispersion states of CNFs was studied. The dielectric
spectroscopy was obtained on a Universal Dielectric Spectrometer BDS 20 from Novocontrol Inc.
using samples with the thickness of ca.1mm. The input voltage (Vrms) was 1V, and the test
frequency range was set between 1Hz and 3MHz. The volume resistivity and surface resistivity
were measured by Keithley 6517A electrometer with 8009 fixture at 50 V. The 8009 fixture has two
large area electrodes for volume resistivity measurement, and ring electrodes for measuring surface
Figure 5.1 Optical microscopic images of dispersion of 3.0wt% CNFs in PEI matrix by Probe
ultrasonication and Bath sonication
5.3 Results and discussion
5.3.1 Electrical and dielectric properties of mono-layer composites
Figure 5.2 Concentration dependence of volume resistivity (A), and surface resistivity of mono-
layer PEI/CNF composites (B)
0 1 2 3 4 5 6 7
Concentration of CNF(wt%)
-1 0 1 2 3 4 5 6 7 810
Concentration of CNF (wt%)
0 1 2 3 4 5 6 7
Concentration of CNF(wt%)
-1 0 1 2 3 4 5 6 7 810
Concentration of CNF (wt%)
First, we will discuss the microstructures and static dissipation property of the two components
composing the bi-layer composites: N-Composite and U-Composite.
In U-Composite, the CNFs treated by high power probe ultrasonication nearly uniformly
disperse in PEI matrix with no obvious CNF agglomerates observed, while in N-Composite, after
bath sonication, the many CNFs exist in PEI matrix in the form of CNF agglomerates (Fig. 5.1).
These two dispersion states affect both volume and surface resistivity significantly. The non-
uniform dispersion of CNFs in N-Composites leads to lower volumetric percolation threshold
between 0.1wt% and 0.5wt%, in contrast with the percolation threshold between 1.0wt%and 2.0wt%
in U-Composites with uniformly dispersed CNFs (Fig. 5.2A). In our study, the volume resistivity
measured by Keithley 8009 fixture is mainly the resistivity along thickness direction. Owing to the
existence of obvious CNF agglomerates in N-Composite, the formation of conductive network
along thickness direction is much easier than uniformly dispersed CNFs in U-Composite.
Compared with volume resistivity, the modification of surface resistivity needs much higher
loading of CNFs (Fig. 5.2B). For the U-Composite, more than 3wt% CNFs is needed to reduce the
surface resistivity to below10-12
ohm/Sq, while for the N-Composite, even at 7wt% loading, the
surface resistivity is only slightly lower than 10-12
Here, we did not study higher loading levels. In polymer nanocomposites, the low filler loading
is usually expected, not only because of the cost issue, but also the ease of processing and
dispersion as well as avoiding generating more structural flaws. Different from volume resistivity,
N-Composite has dramatically higher percolation threshold of surface resistivity than U-Composite.
This is because along the surface direction, the distance between CNF agglomerates is larger, the
conductivity network is more difficult to form compared with uniformly dispersed CNFs obtained
by probe ultrasonication treatment (Fig. 5.1).
Figure 5.3 Relative permittivity of mono-layer composites
U-Composite (1wt% CNFs)
U-Composite (4wt% CNFs)
U-Composite (7wt% CNFs)
N-Composite (1wt% CNFs)
N-Composite (4wt% CNFs)
N-Composite (7wt% CNFs)
U-Composite (1wt% CNFs)
U-Composite (4wt% CNFs)
U-Composite (7wt% CNFs)
N-Composite (1wt% CNFs)
N-Composite (4wt% CNFs)
N-Composite (7wt% CNFs)
As mentioned above, the relative permittivity also significantly impacts on the rate of charge
dissipation of the composites. For an ideal static dissipative material, both low permittivity and
resistivity are expected.
Here, from the permittivity results shown in Fig. 5.3, we found that for U-Composite, although
its surface resistivity has been dramatically decreased with the addition of 4.0wt% probe
ultrasonication treated CNFs, its relative permittivity also exhibits an undesirable increase to ca. 400
at lower frequency range. With 7wt% CNFs loading, the relative permittivity increases further, even
over 2000. While for N-Composite, their relative permittivity just slowly increases and is still below
10 with 7wt%CNF loading. It is reasonable to conclude that the dispersion state of CNFs in two
composites accounts for such a difference in relative permittivity.
In two-phase systems like PEI/CNF composites in our study, the interfacial polarization is the
main mechanism for the dielectric relaxation. Basically, the charge accumulation at the interface
between conductive fillers and polymer matrix leads to the improved dielectric constant (relative
. In U-Composite, due to the nearly uniform dispersion of CNFs, the accumulated
charges can be well retained at the interface when the electric field is applied to the composites.
However, in the N-Composite, a greater number of CNF agglomerates exist. Compared with well
dispersed CNFs, the CNF agglomerates could easily dissipate the accumulated charges owing to
excessive conductive paths within CNF agglomerates. Thus, we can observe the dramatic increase
of relative permittivity in U-Composite. Similar effect of the structure evolution of conductive
network on relative permittivity (dielectric constant) has also been found in PVDF/CNF composites
. In the study of PVDF/graphite nanoplatelet composites, the high relative permittivity has also
been considered as a result of homogenous dispersion of graphite nanoplatelet 
5.3.2 ESD protection performance of mono-layer composites
Based on the above results of electrical resistivity and relative permittivity, the charge
dissipation rate was summarized in Table 5.1. Since the static dissipation property is determined
simultaneously by electrical resistivity and relative permittivity. According to equation (5.1), we
take PEI/7.0wt% CNF composites with the lowest electrical resistivity as an example (both U-
Composite and N-Composite). First we substitute ρ with volume resistivity (this means the static
charges dissipate through the bulk composites), because both U-Composite and U-Composite have
similar volume resistivity around 107
ohm*cm at this concentration, the static dissipation property
will be mainly controlled by relative permittivity. Consequently, for U-Composite with high relative
permittivity (2002.8), its volumetric time constant (ηvol) is much higher than that of N-Composite
showing low relative permittivity (6.5), as given in Table 5.1, indicating poorer volumetric static
dissipation capability for U-Composite. Here, we use the low frequency (1.56Hz) relative
permittivity to do the calculation, because both electrical resistivities were measured in DC field.
Table 5.1 Charge dissipation rate for mono-layer composites
(7.0 wt %)
(ohm/sq) ηvol ηsur
N-Composite 6.5 1.5E7 4.0E10 9.6E7 2.6E11
U-Composite 2002.8 2.7E7 2.3E5 5.3E10 4.7E8
N-Comp/U-Comp 3.2E-3 0.6 1.7E5 1.8E-3 5.5E2
On the other hand, if we substitute ρ with surface resistivity (static charges dissipate in the
vicinity of surface area), the calculated surface time constant (ηsur) will be obtained, and we can also
find that due to the high relative permittivity of U-Composite, the surface static dissipation
capability is also restrained. The difference of ηsur between the two composites (N-Composite/U-
Composite= 5.5E2) is not as large as that of surface resistivity (N-Composite/U-Composite= 1.7E5).
According to these two aspects, it seems that with simply increasing the CNF loading, there is
potential to improve static dissipation performance of composites at low loading of CNFs.
5.3.3 Bi-layer composites with improved ESD protection performance at low loadings
According to the discussion in the last section, we found that in mono-layer PEI/CNF
composites, it is difficult to simultaneously realize low surface resistivity, low volume resistivity
and low relative permittivity at a low loading level. For N-Composite, owing to its remarkable
anisotropy, a great amount of CNFs are needed to decrease both surface and volume resistivity to a
desirable level, while, for U-Composite, the dramatically increased relative permittivity by larger
amounts of uniformly dispersed CNFs is a barrier for superior ESD protection.
Design of bi-layer composite and its relative permittivity
In order to combine the low surface resistivity with low relative permittivity as well as low
loading of CNFs, we designed a type of bi-layer-structured PEI/CNF composites which could
realize these three aspects simultaneously, through a facile solution method, in which, we design a
bi-layer structure by coating a thin layer of U-Composite with low surface resistivity (U-Composite
Layer) on the surface of bulk N-composites with high surface resistivity and low permittivity (N-
Composite Layer). The static dissipation surface is indicated in Fig. 5.4. The interface region in this
bi-layer structure is the key of the design, which is equivalent to a thin membrane composed of two
composite layers with equal thickness as shown in Fig. 5.4.
Figure 5.4 Sketches of bi-layer-structured composites and microstructures of each layer and
In a bi-layer membrane system, its relative permittivity is given by 
𝜀1∗ + (1 − ∅)
where,𝜀1∗ and 𝜀2
∗ are the complex permittivity of two layers. ∅ =𝐷1
𝐷1+𝐷2, D1 and D2 are the thickness
of each layer. Obviously, according to this equation, the real permittivity of the bi-layer membrane
(the real part of 𝜀∗ ) is lower than maximum of the real part of 𝜀1∗ and 𝜀2
Through adjusting the bi-layer structure, we effectively decreased the relative permittivity,
while maintaining both low surface resistivity and low volume resistivity. Figure 5.5 gives the
relative permittivity of two bi-layer structured composites and their mono-layer composites.
Obviously, the relative permittivity of bi-layer structured composite with DU : DN =1:1(8.4) is very
U-Composite layer: Low surface resistivity High permittivity
N-Composite Layer: High surface resistivity Low permittivity
Static dissipative Surface
close to mono-layer N-Composite with 0.5wt% CNFs (3.0), and much lower far than that of
constituent mono-layer U-Composite with 7.0wt% CNFs (2002.8). This is coincident with the
calculated value of 6.3 at 1.56Hz by equation (2). Furthermore, in the case of DU : DN=1:60, the
relative permittivity of bi-layer composite is hardly affected by U-Composite with 7.0wt% CNFs,
and just slightly higher than that of N-Composite with 0.5wt%CNFs (3.0). Thus, according to both
experimental and theoretical results, this bi-layer structure is essentially efficient to yield low
permittivity of PEI/CNF composites.
Single layer U-Composite (7wt% CNFs)
Single layer N-Composite (0.5wt% CNFs)
Bilayer Composite (DU:D
Bilayer Composite (DU:D
Figure 5.5 Relative permittivity of mono-layer composites and bi-layer structured composites
Electrical resistivity of the bi-layer composites
The results of surface resistivity of static dissipative surface as indicated in the bi-layer
structured composite composed of N-Composite layer and U-Composite layer are given in Figure
5.6A (right). In this bi-layer structure, the surface resistivity of static dissipative surface with 0.5wt%
or more CNFs, is as low as ca.106 ohm/sq, in contrast with the surface resistivity of ca. 10
in mono-layer N-Composite with the same amount of CNFs. However, when the concentration of
CNFs in N-Composite layer is 0.1wt%, the measured surface resistivity does not show an obvious
drop, which is similar to the situation of Pure PEI as shown in Fig. 5.6A (left). Taking the volume
resistivity results given in Fig. 5.2 into consideration, the percolation threshold of volume resistivity
of N-composites is between 0.1wt% and 0.5wt%, thus, the result shown in Fig. 5.6 may indicate
that the low volume resistivity of N-Composites layer also contributes to the dramatically reduced
surface resistivity of the static dissipative surface in Fig.6 (right).
Based on this result, we may understand the decrease of measured surface resistivity of static
dissipative surface in the bi-layer structured composite (Fig. 5.6B and 5.6C) by current detouring.
In the mono-layer U-Composite with low surface resistivity as indicated in Fig. 5.6C, the current
directly flows across the materials surface between two electrodes, while for the bi-layer structure
shown in Fig. 5.6B, the volume conduction could help detour the current from the material surface
to the interface region in bi-layer structure as shown in Fig. 5.6B. Because the interface region is
partially composed of U-Composite with low surface resistivity, it provides significant conductivity
to allow the current passing through. Thus, in such a bi-layer–structured composite, a complete
circuit between two points on the surface can form as shown in Fig. 5.6B.
Figure 5.6 Comparison of surface resistivity between bi-layer- structured (U+N) composites with
different concentration of CNFs in layer N and mono-layer N-composites (A). Current flow passing
through two points on the surface: (B) Bi-layer structured composites; (C) Mono-layer U-
Composite with low surface resistivity
That is why the bi-layer structured composites have comparable low surface resistivity (ca.
106ohm/sq) to mono-layer composites (ca.10
Furthermore, the lower relative permittivity of the interface region in the bi-layer structure has
been proved by both experimental results (Fig. 5.5) and theoretical calculation using equation (2),
thus, the charge carriers flowing across any two points on the surface in bi-layer structured
Mono- layer Composite U
PEI 0.1wt% 0.5wt% 1.0wt% 2.0wt%
Concentration of CNFs in Layer N
Single Layer N-Composite
Bilayer structured Composite (U+N)Bilayer structured Composite (U+N)
U-Composite Layer with 7wt%CNFs: Low surface resistivity, high permittivity
Static dissipation surface
N-Composite Layer with varying CNF concentration: High surface resistivity, low permittivity.
composites will be less constrained, compared with mono-layer composites. Thus, it can help to
reduce the time constant for static charge dissipation by simultaneously decreasing surface
conductivity and relative permittivity.
5.3.4 Comparison of ESD between bi-layer and mono-layer composites
According to Fig. 5.6, the bi-layer structured composites composed of U-Composite
(7.0wt%CNFs) and N-Composites (0.5wt%CNFs) can have surface resistivity of ca. 106 ohm/sq,
only one order lower than the U-Composite (7wt%CNFs) (ca.105ohm/sq). In our study, we have
also studied the bi-layer composites with DU: DN of 1:60. In this case, not only the low surface
resistivity of the static dissipative surface can be kept, but also 91.4wt% CNFs can be saved
compared with U-Composite (7wt%CNFs). That means, only 0.6wt% CNF is needed in such a bi-
layer composites to realize a low surface resistivity comparable to that of Composites U (7.0wt%
CNFs). This loading level is lower than most of reported carbon nanofiller modified polymer
composites for static dissipation application [11-13].
The static dissipation property of this bi-layer-structured composites with DU : DN of 1:1 and
1:60 is summarized in Table 5.1. Because both volume and surface resistivity were measured in the
DC field, here, a low frequency (1.56Hz) relative permittivity was used to calculate the time
constant in equation 1. Obviously, the increasing portion of layer N from 1:1 to 1:60 did not
increase the measured surface resistivity of static dissipation surface(3.2E6 ohm/sq to2.4E6 ohm/sq),
while relative permittivity slightly decreased (8.4 to3.8). Considering the very low loading of CNFs,
the bi-layer-structured composites with DU: DN of 1:60 is more ideal for static dissipation
application than the bi-layer-structured composites with DU: DN of 1:1. Thus, the comparison
between bi-layer structured composites and mono-layer composites were made based on the bi-
layer-structured composites with DU: DN of 1:60.
First of all, combined with considerably low volume resistivity and low relative permittivity, the
volumetric static dissipation (ηvol) of this bi-layer structured composites is significantly lower than
that of the mono-layer U-Composite (7wt%CNFs)(only ca. 1/1000 ). Secondly, for the calculation
of surface static dissipation time constant (ηsur), as discussed above, the interface region of the bi-
layer structure is the underlying contributor to the measured low surface resistivity, thus, the
relative permittivity of this interface region is used to calculate ηsur. The structure of this interface
region has been introduced in section 3.1, giving an experimental relative permittivity of 8.39 at
1.56Hz for bi-layer structure with two equal-thickness layer. As given in Table 5.2, the bi-layer
structured composites have a time constant for surface static dissipation (ηsur) of 2.28E7 which is
ca.20 times faster than that of mono-layer U-Composite (7.0wt%CNF), indicating dramatically
improved static dissipation property.
The bi-layer PEI/CNF composites were designed to improve ESD protection capability of
PEI/CNF composites. Through the analysis of volumetric static dissipation and surface static
Table 5.2 Comparison of charge dissipation rate for bi-layer and mono-layer composites
(ohm/sq) ηvol ηsur
wt %) 3.0 1.9E7 4.7E13 5.7E7 1.4E14
wt %) 2002.8 2.7E7 2.3E5 5.4E10 4.6E8
DU:DN =1:1 8.4* 1.8E7 3.2E6 1.5E8 2.7E7
DU:DN =1:60 3.8 2.0E7 2.4E6 7.6E7 2.0E7*
*Note: According to Fig.6B, the calculation of ηsur for Bi-layer Composite (DU:DN=1:60)
also used the permittivity of 8.4, since the current actually goes through the interface region
composed of two composite layers with equal thickness (Fig.4).
dissipation, compared with mono-layer PEI/CNF composites, the static dissipation property of this
bi-layer–structured PEI/CNF composites was remarkably improved, while only needing less than
1/10 CNFs used in mono-layer PEI/CNF composites. The mechanism of the bi-layer structured
composites has been briefly introduced, which strongly depends on the interface region in this bi-
layer structure. Furthermore, the tailorable thickness and thickness ratio of the two layers also
provides versatile applications satisfying various geometry requirements, including antistatic
packaging, microelectronics and even some bulk parts.
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Simultaneous Enhancements in Damping and Static Dissipation Capability of
Polyetherimide Composites with Organosilane Surface Modified Graphene
Polyetherimide (PEI) possesses excellent thermal resistance and advanced mechanical
properties, for which it has been used as interior materials in various transportation structures.
Multifunctional PEI nanocomposites, in particular, with high damping properties and satisfactory
static dissipation capability, will be more significant for those applications. In this study,
PEI/graphene nanoplatelet (GNP) nanocomposites were fabricated via solution processing. Effects
of GNP loading and surface silanization of GNPs on damping capacity and static dissipation
properties were studied. The addition of the GNPs effectively increased the storage modulus of PEI,
especially storage modulus at higher temperature (200°C). The silanization of GNPs, on one hand,
improved the dispersion quality; on the other provided strong interfacial bonding with PEI matrix,
which benefited the stress transfer within the composites. The superior damping capacity
enhancements for the resulting nancomposites are: with the loading of 3.0wt% silanized GNP, the
storage modulus of PEI increased approximately 4 times and 200 times at 300C and 200
respectively, and the damping factor is 3 times higher than PEI. In addition, with the addition of
3.0wt% GNPs, both low electrical resistivity (~106ohm*cm) and dielectric constant (~7) were
realized, corresponding to excellent volume static dissipation capability, however, silanization
resulting in good interfacial bonding did not cause invisible impact on the electrical and dielectric
properties of the nanocomposites. The superior damping capacity and static dissipation property
make them suitable for intra-structure materials for airplane and transportation.
Polyetherimide (PEI) belongs to high performance polyimide family, which has better
processibility because of ether groups on PEI backbones. In recent years, it has gained popularity in
microelectronic, transportation, as well as applications in extreme conditions requiring both
mechanical and thermal stabilities at high temperature. In particular, PEI and its composites are
considered as excellent interior-structure materials for airplanes and ground transportation, which
requires sufficiently high damping capability and satisfactory static dissipation properties.
Graphitic nanofillers, including carbon nanofibers (CNF), carbon nanotubes (CNT) and
graphene nanoplateles (GNP), have been of great interest in the development of functional and high
performance polymer nanocomposites, due to their outstanding mechanical performances, high
thermal stability as well as superior physical properties [1-14]. CNFs [4-5, 7-9, 13] and CNTs [10-
11] have been successfully incorporated into PEI to improve the mechanical, thermal and physical
properties. However, there is no reported work giving exclusive attention to both damping and static
dissipation properties of PEI nanocomposites.
As a reinforcement filler/modifier, GNPs with 2-D nanostructures show some unique
advantages compared with other 1-D graphitic nanofillers (CNF and CNT). At a point of view of
fabrication, the 2-D GNPs can be produced not only by conventional routes used in fabricating
CNFs and CNTs, such as chemical vapor deposition (CVD) and arc discharging methods [15-16],
but also by really convenient approaches like mechanical milling  and graphene intercalation
, as a consequence of their special 2-D structure and abundant mineral graphene in nature. On
the other hand, GNPs can provide multi-functionalities for polymer materials. For example, the 2-D
structures possess greater potential in improving gas permeability of polymeric materials , and
the larger surface area of GNPs increases the interface area between polymer matrix and fillers,
which could benefit interfacial polarization and consequently improve the dielectric capacity of
polymer nanocomposites satisfying the needs for flexible and low weight energy storage materials.
Due to the high electrical conductivity, GNP has been frequently used to improve electrical
properties of polymeric materials, the conductivity level of which could reach as high as 10-4
S/cm , thus GNP is an ideal modifier for static dissipation if it does not cause dramatically
high dielectric constant at the same time [7,12].
As for the damping property, which reflects the ability of materials to convert
mechanical/acoustic vibration into heat energy, the most successful method without obviously
compromising mechanical properties was to fabricate sandwich structure polymer composites
containing CNT films [2, 20]. Through enhanced CNT-CNT friction and CNT-polymer (polymeric
adhesive) friction in the CNT films, the high damping capability has been realized. In another study,
continuous long CNTs were also applied to improve the damping capability of elastomers .
However, for these approaches, special material structures, that is, sandwich structure and long
CNT lengths, are required, which brings difficulties in manufacturing and applications of bulk
polymer nanocomposites, compared with conventional melt and solution processing techniques.
GNPs have a 2-D multilayer structure, in which the unique frictional interlayer sliding normally
exists due to the fact that the graphene layers within GNPs are only bonded with Van der Waals
forces. The interlayer sliding, a characteristic feature of GNP, makes GNP an ideal solid lubricant
that could effectively improve the wear resistance and friction properties of polymers . Thus,
GNPs can have great potential for improving damping properties of polymers, i.e. it is reasonable
to speculate that GNPs as a reinforcement filler can greatly contribute to superior damping capacity
enhancement for matrix materials. However, the research on improving damping properties of
polymers through addition of GNPs has rarely reported.
Hence, in this work, GNPs were chosen for improving both damping capability and static
dissipation properties of PEI, and a silane type surfactant was applied to treat GNPs in order to
improve dispersion and interfacial bonding. Dynamic mechanical analysis (DMA) was used for
evaluating damping capacity of PEI/GNP nanocomposites. The effects of the amount of silane
surfactant as well as the dependence on testing strain were studied. The static dissipation property
(the ability to transport charges) was investigated by measuring the electrical conductivity and
dielectric constant (the ability to localize and store charges). The results indicated that with the
loading of silanized GNPs, the storage modulus increased by 4 times and 200times, respectively at
30°C and 200 °C, while the damping factor of the PEI/GNP composite was 3 times of pure PEI. On
the other hand, the GNPs also dramatically improved the static dissipation property of PEI, which
implies that GNP is superior to CNF .
Polyetheimide (ULTEM1010, Sabic Innovative Plastic, Inc.) and graphene nanoplatelete
(xGnP-5, XG science, Inc) were selected as the matrix materials and reinforcement fillers. The
graphene nanoplaltelet was exfoliated GNPs with a thickness of ~5nm and diameter of ~5μm, as
shown in Fig. 6.1. The technique grade [3- (2-Aminoethylamino) propy]trimethoxysilane (Scheme
6.1a) liquid was the organosilane surfactant purchased from Sigma-Aldrich, Inc.
Figure 6.1 TEM image of ultrasonication treated GNPs used in PEI/GNP nanocomposites
6.2.2 Silanization of graphene nanoplatelet
The GNPs used in this study were exfoliated via an acid intercalation approach, which is a
highly oxidative process and could create abundant chemically active sites on the particle surface,
typically, hydroxyl and carbonyl groups. This is further proved by FTIR spectrum of pristine GNPs
(P-GNP) (Fig. 6.2) in section 6.3.1. Thus, before further surface modification of GNPs, there is no
need to carry out oxidization procedure usually used for graphitic carbon nanofillers. Before surface
modification, 1 gram P-GNPs was dispersed in 100 ml ethanol (95%) by high power ultrasonication
for 1 hr (Branson 450, amplitude 20%). The resultanthomogeneous GNP/ethanol suspension was
removed to a 3 neck flask with a water condenser. The flask was heated up to about 120°C which is
sufficiently high to boil the ethanol. The liquid silane
Scheme 6.1 (a) [3- (2-Aminoethylamino) propy]trimethoxysilane and (b) Silanization of GNPs
surfactant was dropwise injected into the sealed 3 neck flask. After 5hrs reaction, the GNP/ethanol
suspension was cooled down to room temperature, and the silanized GNPs (S-GNP) was rinsed and
collected via filtration, and then dried at 70 °C for further use. Two silane / GNP ratios were
selected: 4ml / 1g and 6ml /1g.
During this process, the silane surfactant reacted with the active groups, such as hydroxyl
groups, according to scheme 1b. The small amount of water in ethanol helped the hydrolyzation of
–OCH3 groups in silane molecules. The resulting –SiOH groups could react with the hydroxyl and
carbonyl groups on GNP surfaces and formed strong covalent bonds (-Si-O-C), as well as adjacent
R: - (CH2)3NH(CH2)2NH2
R’ : - CH3
–SiOH groups. At the same time, due to the existence of both hydrogen and hydrogen acceptor,
some minor hydrogen bonding might also forms. On the other hand, the rest part of silane
molecules is compatible with polymer matrix, and consequently improve the interfacial bonding.
6.2.3 Preparation of PEI/GNP nanocomposites
The solution processing method was applied to fabricate PEI/GNP composites. Both P-
GNPs and S-GNPs were dispersed in methylene chloride (CH2Cl2), respectively, by high power
ultrasonication (Brason 450, amplitude 20%) for 1hrs. After adding PEI powder to the suspension,
spin mixing was used to dissolve PEI and mix PEI and GNPs at a speed of ~ 40-50 rpm for 12 hrs.
The resulting PEI/GNP solutions were treated by low power bath sonication (Branson 1510) for
another 2 hrs. The film samples of PEI/GNPs composites with a thickness of ~30um were obtained
via cast coating. Since the viscosity of polymer solution has remarkable impact on the dispersion of
nanofillers, in order to maintain the similar dispersion quality, the polymer to solvent ratio was
fixed at 3g /10ml. For scanning electron microscopy, electrical and dielectric tests, the 1mm thick
panel samples were prepared by compression molding of film sample at 600°F. Table 6.1 gives
the PEI/GNP nanocomposites studied in this work.
6.2.4 Fourier transform infrared spectroscopy
In order to confirm the success of silanization, Fourier transform infrared spectroscopy
(FTIR) (Thermo Nicolet 6700) was applied to both P-GNP and S-GNP. Approximate 0.02 g GNPs
was mixed with 1g KBr powder via mechanical grinding. The circular panel sample was obtained
by compressing fine mixture powder under high pressure. The FTIR spectra are given in Fig. 6.2.
Table 6.1 List of PEI/GNP nanocomposites in this study
Designation Concentration of GNPs
Amount of silane surfactant
PEI 0 0
PEI/p-0.5GNP 0.5 0
PEI/4s-0.5GNP 0.5 4
PEI/6s-0.5GNP 0.5 6
PEI/p-3.0GNP 3.0 0
PEI/4s-3.0GNP 3.0 4
PEI/6s-3.0GNP 3.0 6
p – Pristine; s- Silane treated; 4 -4ml silane/g GNP; 6 – 6ml silane /g GNP
6.2.5 Dispersion and interface analysis
Dispersion analysis was carried out by both studying the stability of GNPs in dilute
PEI/CH2Cl2 solution, as shown in Fig. 6.3, and scanning electron microscope (SEM) (Fig. 6.4).
0.005 g GNPs was directly dispersed in dilute PEI/ CH2Cl2 solution (0.2g/5ml) via high power
ultrasonication for 10mins to obtain homogeneous PEI/GNP suspension. Then the time dependence
of solution stability was recorded up to 2 months. At the same time, the cross-section of cryo-
fractured PEI/GNPs nanocomposites was observed by SEM (FEI 200) to study both dispersion
quality and interface morphology.
6.2.6 Dynamic mechanical analysis
The 10mm (length) ×8mm (width) samples were cut from 30μm thick PEI/GNP
nanocomposites films for tension mode DMA testing (Tritec 2000). The scan frequency was set at
1Hz. With a ramp rate of 3°C/min, the storage modulus (E’) and damping factor (tan(δ)) were
recorded in the temperature range of 30°C ~ 300°C. In particular, the dynamic mechanical
properties under different testing strain were studied. All chosen testing strain fell into linear elastic
region determined by strain sweep at room temperature. For each nanocomposite, at least 5 samples
were test at each test strain.
6.2.7 Electrical and dielectric properties
Volume DC resistivity of PEI/GNP nanocomposites was measured using Keithley 6517A
electrometer equipped with 8009 test chamber for high resistance materials; while the dielectric
constant was studied on a Universal dielectric spectrometer BDS 20 from 0.01 Hz to 106Hz.
6.2.8 Thermal analysis
Thermal stability of PEI and PEI/GNP nanocomposites with and without surface modification
was examined by thermal gravimetric analyzer (TGA, TA Q600) at a ramp rate of 20°C/min up to
800 0C in N2 atmosphere.
6.3 Results and discussion
500 1000 1500 2000 2500 3000 3500
Figure 6.2 FTIR spectrums for pristine GNPs (p-GNP) and 4ml silane treated GNPs (4s-GNP)
In Fig. 6.2, the FTIR spectrum of P-GNPs shows a broad peak around 3400cm-1
to –OH groups, suggesting there are plenty of active sites for silanization of P-GNPs. Typically,
The previous work in our group revealed that , for silane surfactant, there should be two weak
peaks around 800cm-1
, due to the bending of Si-O-CH3 and stretching of -Si-O groups.
As a result of silanization, the peak at 800cm-1
would disappear, while the 1000-1
cm peak became
really broad. However, for P-GNP, there is a strong peak around 1000cm-1
from GNP itself, which
could overlap the peak of stretching of -Si-O groups in S-GNP. Thus, we cannot testify the
silanization of GNPs by the broadened peak at 1000cm-1
. Notice that there are four peaks around
(inserted figure in Fig. 6.2), including a broad peak around 900 cm-1
GNP , which do not exist in P-GNP, indicating the formation of coating layer of silane molecules on
GNP surface in S-GNP. The broader and stronger peaks at 900 cm-1
were ascribed to chemical
bonds of Si-O-C (GNP), while the ones at 800cm-1
might reflect the existence of some unreacted Si-
O-CH3 groups on GNPs. In addition, for the peak at 1000cm-1
, it is obvious that its peak strength
decreased after silanization. Because we used the same amount of GNPs as well as the same
GNP/KBr ratio in FTIR study, such a decrease may also indicate the formation of a coating layer,
which disturbed the FITR absorption of GNPs. At last, the 3400cm-1
peak was hardly affected,
although most –OH groups have been reacted, and this is because the existence of amino groups on
GNP surface, whose characteristic peak is also around 3400cm-1
6.3.2 Stability of PEI/GNP solution
Besides FTIR analysis, the investigation of stability of PEI/GNP solution can also prove the
successful surface modification of GNPs by silanization. Although for both P-GNPs and s-GNPs,
they can homogeneously disperse in PEI/CH2Cl2 dilute solution after the ultrasoncation treatment,
the difference between them became notable shortly in a few hours. Specifically, the P-GNP
without surface modification started to aggregate on the top of the dilute solution, and within 24 hrs,
it was hardly to observe p-GNP particles in the PEI solution (Fig. 6.3). For both s-GNPs, the s-
GNPs readily stayed in the PEI solution, even after 2 months, the suspensions still looked rather
homogeneous, suggesting that first, the silanization was successful; and second, the R groups in
scheme 1 had strong interaction with PEI chains, due to the existence of polar amino groups. It
should be pointed out here, that our early work has found that the high power ultrasonication could
break down covalent bonds in CNFs and shortened the length . In this part of the study, we used
the similar high power ultrasonication to disperse GNPs before observing their stability. Therefore,
the superior stability of s-GNPs in dilute PEI solution further revealed that silane molecules were
mostly chemically bonded to the GNP surface via the reactions given in scheme 6.1. At the same
time, unlike p-GNPs who aggregated above the solution surface, it was observed that s-GNPs
inclined to fall down to the bottom, suggesting the significant changes of surface properties of s-
GNPs which consequently affected interaction of GNPs with PEI as well as mass density of s-GNPs.
Figure 6.3 Stability of GNPs in dilute PEI/CH2Cl2 solution after two month
p-GNP 4s-GNP 6s-GNP
Agglomerates of p-GNPs
Poor interface (gap)
B Strong interface
Figure 6.4 SEM images of cryofracture surfaces of (A) PEI/p-3.0GNP, (B) PEI/4s-3.0GNP
nanocomposites (low magnification) and (C) PEI/4s-3.0GNP nanocomposites (high magnification)
6.3.3 Morphology of PEI/GNP nanocomposites
The SEM images in Fig. 6.4 provide insightful morphological information for both p-GNPs and
s-GNPs filled PEI nanocomposites. Thus, only the images of PEI/p-3.0GNP and PEI/4s-3.0GNP
nanocomposites are given in Fig. 6.4. Generally, the dispersion technique used in this study could
lead to a good dispersion quality of both GNPs in PEI matrix. However, for p-GNPs filled PEI
nanocomposites (Fig. 6.4A), big agglomerates still exist. Moreover, the poor interfacial bonding
Uniformly dispersed s-GNP with good
between p-GNPs and PEI matrix created obvious gaps, due to lack of effective interface
modification. At the same time, more voids were also observed on the fracture surface, which
should be a result of fall-off of p-GNPs during cryo-fracturing. Both dispersion and interface have
been effectively improved by silanization of GNPs, as shown in Fig. 6.4B. The s-GNPs uniformly
dispersed and distributed in PEI matrix with much better interface conditions. Neither the voids nor
the gaps between s-GNPs and PEI could be observed, indicating the interface bonding has sufficient
mechanical strength to resist external forces when fracturing nanocomposites for SEM studies.
To sum up, from the FTIR analysis, stability test and SEM images, it can be concluded that via
the silanization procedure shown in scheme 1, the silane molecules could be successfully bonded
onto GNP surfaces, and consequently improve the dispersion and interface bonding, due to the
effective interaction between silane surfactant with PEI matrix.
6.3.4 Dynamic mechanical /damping properties
The studies on the storage modulus of PEI/GNP nanocomposites at different testing strains were
carried out at four representative temperatures from room temperature to adjacent glass transition:
30°C, 100°C, 150°C and 200°C, since the glass transition temperature of PEI is around 220°C, and
the dynamical mechanical properties at higher temperature are not the interest of this work.
Obviously, pure PEI has the lowest storage modulus (Fig. 6.5). In our testing range, the storage
modulus of PEI slightly decreased with strain to a level-off value, and did not show significant test
strain dependence, probably due to the fact that all testing strains are in linear elastic region. Similar
results have also been found in PEI/GNP nanocomposites in this study. In contrast, the temperature
dependence is evident for pure PEI. When temperature is increased to 200°C, there is a dramatic
drop of storage modulus, because of getting close to glass transition temperature. According to Fig.
6.5 and Fig. 6.6, the addition of GNPs as well as surface modification could restrain this drop, that
is, the storage modulus of the composites could maintain a high level similar to that at room
temperature, which could significantly benefit the high temperature applications. However, for
PEI/GNP nanocomposites, the nano-reinforcement and surface silanziation of GNPs play different
roles at low and high loading levels, which will be discussed later.
Figure 6.5 Storage modulus of 0.5 % GNPs filled PEI nanocomposites: effect of temperature and
0.0 0.2 0.4 0.6 0.8 1.0
0.0 0.2 0.4 0.6 0.8 1.0
0.0 0.2 0.4 0.6 0.8 1.0
0.0 0.2 0.4 0.6 0.8 1.0
Figure 6.6 Storage modulus of 3.0% GNPs filled PEI nanocomposites: effect of temperature and
Figure 6.7 Effect of addition of GNPs as well as silanization on the improvement of storage
modulus of PEI/GNP nanocomposites at 0.01% strain
When only 0.5wt% p-GNP was added to the PEI matrix, the storage modulus at 30°C showed
an approximate 100% increase, however, the much more significant improvement in storage
20 40 60 80 100 120 140 160 180 200 220
Materials Slope of linear regression line
PEI/ P-3.0GNP -0.001
PEI/ 6S-3.0GNP -0.001
20 40 60 80 100 120 140 160 180 200 220
Materials Slope of linear regression line
PEI/ P-0.5GNP -0.004
PEI/ 4S-0.5GNP -0.003
PEI/ 6S-0.5GNP -0.002
0.0 0.2 0.4 0.6 0.8 1.0
0.0 0.2 0.4 0.6 0.8 1.0
modulus was found at 200°C, which reduced the differences of the storage modulus between room
temperature and high temperature, suggesting potential high temperature applications of PEI/GNP
nanocomposites. The silanization also showed remarkable contribution to the storage modulus.
With the increase of the amount of silane from 0ml to 6ml per gram GNPs, the storage modulus at
all three temperatures gradually increased. In particular, for PEI/6s-0.5GNP nanocomposite with
highest silane/GNP ratio, its storage modulus at 200°C is almost close to that at 30°C. Compared
with pure PEI, the 0.5wt% s-GNP with a silane/GNP ratio of 6ml/g could lead to 2 times and 100
times increase in storage modulus at 30°C and 200°C, respectively. Further improvement in storage
modulus can be realized by increasing GNP loading to 3.0wt% (Fig. 6.6). Specifically, both P-
GNPs and s-GNPs could lead to 4 times at 30°C and 200 times increase at 200°C at 3.0wt% loading
Obviously, both the nano-reinforcement and surface modification of GNPs could increase the
storage modulus of PEI, especially for the high temperature modulus as discussed above. However,
comparing these two groups of nanocomposites with different loading levels, it is easy to find out
that the contributions of nano-reinforcement as well as surface silanization of GNPs strongly
depend on the loading levels. In order to provide more explicit and quantitative information of their
contribution, following analysis was carried out. In storage modulus vs. temperature curves (Fig.
6.7), we did linear regression to the storage modulus at three temperatures, and the slope (no real
physical meaning) of the regression lines was used to evaluate the contributions of both factors. The
smaller the absolute value of the slope is, the better the reinforcement effect is. Here, only the data
obtained at 0.01% strain were given, because of similar results at other strains. For pure PEI, its
slope is -0.01 with the highest absolute value among all the materials studied. This absolute value
was quickly reduced to 0.004 with the addition of 0.5wt% p-GNPs. The surface modification further
decreased this value to 0.002. At the same time, with 3.0wt% p-GNP loading, the bigger reduction
to 0.001 can be obtained, however, at this loading level, the silanzation of GNPs did not show
significant contribution to the enhancement of storage modulus. For PEI/6s-3.0GNP nanocomposite,
the absolute value of the slope is also 0.001.
Figure 6.8 Illustration of two types of GNP networks in polymer matrix (d is the critical inter-
particle distance to form an effective GNP network).
According to the results of Fig. 6.7, we can conclude that the effect of surface modification is
more remarkable at low GNP loading level. In rigid fillers reinforced polymer composites, the
Without surface modification Poor interface bonding
High GNP loading
With surface modification Strong interface bonding
Low GNP loading
GNP Inefficient stress transfer Efficient stress transfer
increased storage modulus of polymers basically results from two mechanisms : (1) the intrinsic
high modulus of the rigid fillers; (2) the formation of network structure of rigid fillers, which could
increase the load capacity of polymer matrix. The first reason can briefly explain why only 0.5wt%
p-GNPs (low loading level) could obviously improve the storage modulus of PEI in a wide
temperature range; while the understanding of the contributions of silanization needs to understand
the network structures of GNPs in PEI matrix, with the help of schematic illustration in Fig. 6.8.
Basically, the formation of networks requires the inter-particle distance among rigid fillers smaller
than a critical value, d (Fig.8). For electrical properties, when the inter-particle distance is smaller
than d, electron hopping takes places, and the materials became electrically conductive; for
mechanical properties, the stress fields around filler particles starts to interact with each other, and
consequently affect the mechanical properties. This concept has been frequently used to explain the
elastomer and rigid particle toughening mechanism [23-24]. Usually, both concentration of fillers
and interfacial bonding have essential influences on this critical distance (d) and formation of
At high loading levels of GNPs, such as 3.0wt%, owing to great number of GNP particles, the
inter-particle distance is rather small, and easily satisfies the critical value (d). Thus, no matter if the
interface condition is good or not, with or without interface modification, the small interparticle
distance is sufficient for the construction of GNP networks, corresponding to TYPE I in Fig. 6.8. At
the same time, the high modulus of GNPs also plays an important role. On the other hand, for the
PEI/GNP nanocomposites with only 0.5wt% GNPs, the large interparticle distance prevents the
formation of TYPE I network, however, the improvement in interfacial bonding via silanization not
only can improve the dispersion (break down the agglomerates, increase the number of individual
GNP particles and decrease the inter-particle distance), but also helps the efficient stress transfer,
which directly enhances the interactions among different stress fields, and equivalently enlarges the
critical value (d) needed to form the GNP network. Thus, the effective interfacial bonding can help
form the TYPE II network at lower GNP loading (Fig. 6.8).
0.00 0.02 0.04 0.06 0.08 0.100.01
Figure 6.9 Damping factor (tan (δ) of PEI/GNP nanocomposites at room temperature
The damping factors (tan(δ)) measured at 0.01% and 0.05% were given in Fig. 9. The
incorporation of p-GNPs in PEI matrix can only slightly increase the tan(δ), even at 3.0wt% loading
of p-GNPs, the increase amplitude of tan(δ) is still smaller than 0.01, from ~ 0.021 to ~ 0.03. In
contrast, the silane treated GNPs (s-GNP) show great damping capabilities, in particular, at high
silane content (6ml/g). Even at 0.5wt % loading level, the resulting nanocomposite, PEI/6s-0.5GNP,
has a high tan(δ), only second to PEI/6s-3.0GNP composite whose tan(δ) reached ~0.07, compared
with ~0.021of pure PEI. This result reveals that, to damping capability, the interface
bonding/surface modification is much more important than the concentration of GNPs, which is not
consistent with storage modulus. The reason for high damping capacity for both PEI/6s-0.5GNP
and PEI/6s-3.0GNP composites can be understood as a result of enhanced frictional motions,
especially the interlayer sliding of GNPs, due to more efficient stress transfer in composites with
strong interfacial bonding. On the other hand, the silanization of GNPs also affected the strain
dependence of tan(δ), the nanocomposites filled with p-GNPs did not show as much improvement
as s-GNPs filled PEI nanocomposites (except for PEI/6s-3.0GNP composites), when the strain
increases to 0.05% from 0.01%. At higher strain, larger deformation is able to strengthen frictional
motions among fillers and polymers. However, owing to the lack of sufficient stress transfer, the
larger deformation of PEI/p-0.5GNP and PEI/p-3.0GNP nanocomposites could not effectively lead
to the enhanced frictional interlayer sliding in GNPs as well as the frictional motion among GNPs
and PEI matrix.
Lastly, both nano-reinforcement of GNPs and improved interfacial bonding via silanization did
not cause any increases in glass transition temperature (Tg) of PEI. Compared with pure PEI, the
variation of Tg for all composites is within ± 2 °C which is considered as an instrumental error. This
fact appears to be good news for the processing of PEI/GNP nanocomposites. For amorphous
polymers like PEI, the melt processing temperature greatly relies on their Tg. PEI has high Tg as
well as high melt viscosity, which makes the processing of PEI and its composites challenging, and
its melt processing temperature is usually recommended as high as above 340°C. The non-
increasing Tg of PEI/GNP nanocomposites could prevent to a certain degree, the further increase of
processing temperature. Unlike storage modulus, more structure factors determine Tg in polymer
nanocomopsite. Usually, in polymer/graphitic nanofiller composites, it is believed that strong
interfacial bonding could increase Tg . However, for PEI/GNP nanocomposites, there are
several reasons possibly responsible for the unchanged Tg, despite of good interface bonding. First
of all, compared with the rigid backbone structures of PEI with poor mobility, the interlayer sliding
in GNP seems much easier, thus, instead of restricting the motion of PEI chain segments, the strong
interfacial bonding would rather induce the friction motions in GNPs; Secondly, the interaction
among PEI and the organosilane (scheme 1) with short R group might not be strong enough.
6.3.5 Static dissipation property
The static dissipation property was studied based on electrical and dielectric properties. The
electrical conductivity indicates the ability of materials to transport electrical charges, while
dielectric constant is a competing factor which represents the ability of materials to
store/accumulate charge carriers. Thus, a time constant τ, the product of electrical resistivity (ρ) and
dielectric constant (ε), could be used for evaluating the static dissipation properties, as given in
equation (6.1) :
τ = ρε (6.1)
The smaller the time constant is, the faster the static dissipated, i.e. the better the dissipation
property is. However, usually, the addition of conductive fillers could also increase the dielectric
constant. Obviously, a balance between electrical resistivity and dielectric constant is needed.
With only 0.5wt% GNPs, the PEI/GNP nanocomposites still act as insulating materials as pure
PEI. For PEI/3.0wt% GNP nanocomposites, the resistivity shows a sharp drop to around
106ohm*cm, indicating the potential for static dissipation, indicating that the percolation threshold
for forming conductivity network should be between 0.5 wt% and 3.0wt%, which is not the subject
of this work. At the same time, a pleasant result is that for both 0.5wt% and 3.0wt% nancomposites,
the surface modification via silanization did not deteriorate the electrical properties. According to
previously reported work, the surface modification of conductive filler by insulating surfactant
could harm the electrical properties of resulting composites by weakening electron transport, which
did not happen to PEI/GNP composites, probably due to the small molecules of silane surfactant
and thin surfactant layer incapable of stopping the motion of electrons.
PEI/P-0.5GNP PEI/6S-0.5GNP PEI/P-3.0GNP PEI/6S-3.0GNP
Figure 6.10 Volume electrical resistivity of PEI/GNP nanocomposites
The dielectric constant for all materials are given in Fig. 6.11, similar to electrical properties, it
is still the concentration of GNPs dominating the dielectric properties, while the silane surfactant
has little effect. The dielectric constant increased with the loading of 0.5wt% GNPs, at
Figure 6.11 Dielectric constant of PEI and PEI/GNP nanocomposites
the same time, just like PEI, the dielectric constant did not show obvious frequency dependency. By
increasing loading to 3.0wt%, the dielectric constant exhibited an increase to over 7 accompanied
with obvious frequency dependence, suggesting an enhanced interfacial polarization mechanism.
Actually, the dramatic improvement in dielectric constant via adding GNPs to polymer matrix has
been frequently reported [25-27], which was typically several orders higher than our results,
probably because of the properties of polymers and processing techniques. However, for antistatic
application, this low dielectric constant is welcomed, thus, our GNP reinforced PEI nanocomposites
have more promising potentials in static dissipation than other GNP reinforced polymer
nanocomposites. For PEI/3.0wt% GNP composites (both PEI/6s-3.0GNP and PEI/p-3.0GNP), the
time constant τ for volume static dissipation is only approximately 2E7 which is ca. 3000 times
faster than PEI/7.0wt% CNF nanocompoosites (5.4E10), and slightly higher than that of the bi-
layer PEI/ 0.6wt% CNF nanocomposites (7.6E7) , indicating the remarkably superiority of GNPs
to CNFs in static dissipation applications.
6.3.6 Thermal stability
It has been frequently reported that the addition of GNPs in polymer matrix could improve the
thermal stability. Thus, in this study, the thermal stability of PEI/GNPs was studied by thermal
gravity analysis (TGA) in N2 atmosphere. According to Fig. 6.12, the reinforcement of 3.0wt%
pristine GNPs in PEI matrix could effectively increase the onset temperature of thermal degradation
(5wt% weight loss) by 16.6 0C. Several mechanisms  have been explained the contribution of
GNPs on the increased thermal stability of polymers: 1) the homogeneously dispersed GNPs acts as
“efficient heat sinks”, which consumed more heat than the matrix and did not allow the
accumulation of heat within the latter, and thereby prevented oxidation at the early stages of
degradation; 2) the homogeneously dispersed GNPs could serve as the mass transfer barriers
(shielding effect) against the volatile pyrolized products; 3) the interfacial polymer phases in the
vicinity of the graphite nanoparticle surfaces are restricted by the bonding from GNPs, and the
energy needed to decomposition would increase, alter the ability of degraded molecules to diffuse
and evaporate. According to Fig.6.4, we can see the agglomerates of GNPs as well as poor
interfacial bonding, thus, it is believed the first mechanism may be the main reason for the 16.6°C
increase. However, for the surface modified GNPs, they did not lead to as much increase as pristine
GNPs did, and only 7.4 °C increase could be observed, despite of uniform dispersion of GNPs as
well as stronger interfacial bonding. It is assumed that this is resulted from the lower thermal
stability of surface coating layer of organo-silane surfactant, suggesting that, rather than dispersion
and interfacial bonding strength, the thermal properties of interface modifier might also be a critical
factor when fabricating polymer/GNP nanocomposites with substantially improved thermal stability.
0 200 400 600 800 10000.4
Figure 6.12 Thermal gravity analysis of PEI and PEI/GNP nanocomposites
The nano-reinforcement by graphene nanoplatelet (GNP) as well as surface modification of
GNPs via silanization simultaneously improved damping properties and static dissipation capability
for PEI. In particular, the storage modulus and damping capacity of PEI were remarkably improved.
The facile silanization showed high efficiency in improving both properties by strengthening the
interfacial bonding. Results also showed that with a proper loading of the silane treated GNPs in
PEI, the volume static dissipation capability of the nanocomposites can be boosted to a very high
level, without being impaired by the existence of the insulating silane layer on GNP surfaces.
Therefore, this study suggests that the resultant PEI nanocomposites with the silanized GNPs are
promising for the applications with high damping capacity and excellent static dissipation property,
as well as high temperature stability, such as airplane/transportation interior structures. Moreover,
this study further proves the great potential of graphene materials as ideal multi-functional
nanofillers for high performance polymer nanocomposites.
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Effect of Non-Covalent Surface Modification via Poly(3,4-
ethylenedioxythiophene)-poly(styrenesulfonate) on Electrical Properties of
Porous Polyetherimide/Carbon Nanotube Nanocomposites
Non-covalent surface modification was applied to carbon nanotube (CNT), in order to fabricate
polyetherimide (PEI) /CNT composites with superior electrical properties. Water soluble and
electrically conductive poly(3,4-ethylenedioxythiophene)-poly(styrenesulfonate) (PEDOT:PSS) was
chosen as the non-covalent surfactant. The results revealed that PEDOT:PSS could efficiently
improve the dispersion of CNT in both aqueous solution and PEI matrix. The resultant PEI
nanocomposites with non-covalently modified CNTs exhibited extraordinary volume and surface
electrical properties. The volume resistivity along sample length direction and surface resistivity
were as low as ca. 8000 ohm·cm and 5×106 ohm·sq, respectively, with only 3.0wt % CNT loading,
lower than 107
ohm·cm and 1011
ohm·sq in 3.0wt% pristine CNT reinforced PEI nanocomposites.
It is believed that the complex conductive network of uniformly dispersed CNTs as well as the
PEDOT:PSS macromolecules which provided extra conductive pathes is the main reason for the
magnificent electrical properties.
Surface modification of nanomaterials plays an important role in developing high performance
polymer nanocomposites. Proper surface modification could substantially improve properties and
functionalities of polymer nanocomposites via not only improving dispersion and distribution
qualities of nanomaterials in polymer matrix, but also enhancing interactions among nanomaterials
and polymers [1-2]. In addition, some surfactants, such as conductive polymers with conjugative
structures, are functional materials with unique physical and chemical properties. Their existence in
the interface between polymer matrix and nanomaterials often lead to extraordinary functionalities
in nanocomposites [3-4].
Graphitic carbon nanofillers( GCN), including carbon nanotube (CNT), carbon nanofiber (CNF)
and graphite nanoplatelet (GNP) exhibit extraordinary mechanical, thermal as well as physical
properties, due to their unique sp2 hybridization of carbon atoms [5-12]. Thus, they are considered
as a group of all-purpose nanomaterials extensively used in high performance and functional
polymer nanocomposites. It has been frequently reported that the addition of GCNs in polymer
matrix could not only amend the low strength, modulus and thermal stability of polymeric materials,
but also succeed in tailoring their electrical, dielectric as well as optical properties for advanced
applications. In order to achieve high performances and functionalities in GCN reinforced polymer
nanocomposites, besides dispersion and distribution control of GCNs in polymer matrix, the surface
modification is also an efficient approach. In particular, due to the sp2 hybridized structures, the
surface modification cannot only be accomplished by conventional covalent chemistry method, but
also non-covalent chemistry method. In covalent surface modification method, reactive groups,
such as -OH and –COOH, etc, are firstly introduced to the GCNs surface, before attaching
surfactant molecules to GCNs, due to the chemically inert sp2 hybridized carbons. Covalent surface
modification is the most frequently used and well understood surface modification method for
GCNs. The covalently bonded surfactant molecules could effectively prevent the aggregation
among GCNs, and lead to strong interfacial bonding between GCNs and polymer matrix. It has
achieved greatest success in improving mechanical, thermal and tribological properties of polymers
[1, 2, 13-16].
However, in order to create sufficient reactive sites on GCNs, the intense oxidization is often
required, which usually leads to conversion of sp2 hybridization to sp
3 hybridization. According to
previous studies on GCNs , it is believed that sp2 hybridization is the fundamental structural
reason for the extraordinary mechanical, thermal and physical properties of GCNs, thus, one
consequence of covalent surface modification is the declined properties and functionalities of GCNs.
Furthermore, the surfactants used in covalent surface modification, such as coupling agents , are
usually lack of functionalities. Thus, although they could successfully improve the dispersion of
GCNs and strengthen interfacial bonding, they hardly have significant contributions to physical
properties of GCN reinforced polymer nanocomposites, such as electrical and dielectric properties.
In contrast, non-covalent surface modification makes up the incompetence of covalent surface
modification in achieving functionality in polymer nanocomposites. Conductive polymers are
typical non-covalent surfactants, because of their conjugative structures. Non-covalent surface
modification by conductive polymers is based on the π-π interactions between conjugative
structures and sp2 hybridized carbons. However, conductive polymers are not compatible / miscible
with other polymers. In order to improve the interfacial bonding, the copolymers of conductive
polymer with compatible polymer blocks are recommended. Non-covalent surface modification has
gained its reputation via its ability to maintain the unique sp2 hybridized carbon structures, and
consequently keep the superior properties of GCNs. In addition, the conjugative structures in
conductive polymers also account for the special physical and chemical properties, thus, the non-
covalent surface modification of GCNs is considered more efficient in functionalizing polymer
nanocomposites, compared with covalent surface modification of GCNs. Previous studies did not
only show the effect of non-covalent surface modification on decreased electrical percolation
threshold and higher level-off conductivity , but also showed its potential in developing polymer
nanocomposites dielectric with high dielectric constant and low dielectric loss . So far the
capability of non-covalent modification to uniformly disperse GCNs has been frequently reported
[20-23], however, the studies on polymer nanocomposites modified via non-covalent surface
modification are still limited [3, 4, 19].
Polyetherimide (PEI) is an engineering thermoplastic with excellent mechanical and thermal
properties, showing great potential in electronics and infrastructures for airplane and ground
transportation [25-26]. Both anticipated applications of PEI require tailorable physical properties, in
particular, electrical properties, to optimize funtionalties and /or avoid electrical discharging for
materials safety. GCNs, including CNT, CNF and GNP have been successfully applied to high
performance PEI nanocomposites, resulting in improvement of mechanical properties, tribological
properties as well as electrical properties of PEI [24-26]. Regarding electrical properties of
PEI/GCN nanocomposites, the low percolation threshold of volume electrical conductivity between
0.3wt% ~ 0.5 wt% for multi-walled CNT has been reported , while the hybrid nano
reinforcement by GNP and CNT showed lower percolation threshold then single nano
reinforcement . However, two existing problems still in the way of developing high efficiency
electrically conductivity PEI/GCN nanocomposites: (1) moderate level-off volume conductivity and
(2) low surface conductivity.
Today, achieving extremely low percolation threshold of volume electrical conductivity in
polymer nanocomposites is not quite an exciting research challenge. Percolation threshold as low as
0.01wt% or even much lower has been frequently reported in many polymeric materials .
However, high level-off volume conductivity at low loading level is still more challenging task,
which is strictly limited by the tunneling conduction mechanism in conductive polymer
nanocomposites with low loadings of GCNs. Conventional approaches to improve level-off volume
conductivity includes aligning and orienting GCNs by external driving forces, such as electrical and
magnetic fields, in polymer matrix to gain high electrical conductivity along the orientation
direction, as well as increasing the loading of GCNs. However, both approaches have their
limitations, including complicated processing procedures, loss of low mass density and flexibility of
polymers as well as high cost and so on. Meanwhile, the study on surface conductivity has not
gained as much attention as volume conductivity, although it is a critical parameter to evaluate the
static dissipation capability which is important to applications such as airplane, ground
transportation as well as electronics, etc. Our previous study on CNF modified PEI nanocomposites
 has already revealed that the concentration and dispersion dependences of surface conductivity
were inconsistency with those of volume conductivity. In particular, in order to improve the surface
conductivity to a satisfactory level, more CNFs were needed. Thus, how to simultaneously improve
both surface and volume conductivities in polymer nanocomposites at a low loading level is still a
challenging but valuable work.
The main objective of this study is to significantly improve both the level-off volume
conductivity and surface conductivity of PEI/GCN composites at a low loading level via non-
covalent surface modification of CNT with water soluble poly(3,4-ethylenedioxythiophene)-
poly(styrenesulfonate) (PEDT: PSS). The emulsion of chloroform/deionized water was applied as
solvent to disperse non-covalently treated CNT in PEI matrix. The water phase could not only help
the dispersion, but also result in porous structures leading to the low mass density of
nanocomposites. The resultant porous PEI/CNT nanocomposites were examined by electrical
properties and dielectric properties.
Polytherimide (ULTEM 1000) (PEI) was provided by Sabic Innovative Plastic Inc; Multi-
walled Carbon Nanotubes (Grade: DM, 99.9% carbon) (CNT) was purchased from Catalytic
Materials LLC; Aqueous solution of highly conductive Poly (3, 4-ethylenedioxythiophene)-poly
P) (PEDOT:PSS) with maximum conductivity of 1 S/cm was
purchased from Heraeus Conductive Polymers Division. The chemical structure of PEDOT: PSS is
given in Scheme 7.1. In order to precisely control the ratio of CNT to PEDOT: PSS, the aqueous
solution of PEDOT:PSS was dried at 700C to obtain weighable solid PEDOT:PSS.
Scheme 7.1 Chemical Structure of Poly (3, 4-ethylenedioxythiophene)-Poly (styrenesulfonate)
7.2.2 Surface modification of CNT
According to reported work on non-covalent surface modification of GCNs , the surface
modification of CNT could be accomplished by dispersing pristine CNT in aqueous solution of
PEDOT: PSS via high power ultrasonication. In this study, 0.1 g CNT was dispersed in 15ml
aqueous solution containing 0.125 g PEOT: PSS by ultrasonifier (Brason 450) for 1hour in ice
water bath. The power intensity was set at ca. 80 watts. The surface modified CNT was dried at 70
0C for preparation of nanocomposites and examination of surface modification. P-CNT and S-CNT
represent pristine CNT and surface modified CNT in this study.
7.2.3 Preparation of PEI/PEDOT:PSS/CNTnanocomposites
The emulsion of chloroform/ water (v/v =9/1) was prepared by high power ultrasonication for
10 mins. After adding pristine and surface modified CNTs to the emulsion, the high power
ultrasonication was applied to the suspension for 20mins to disperse CNTs. Because of the
existence of water soluble PEDOT: PSS on surface modified CNTs, they can be easily dispersed
into the water phase of the emulsion. PEI powder was added to the suspension of CNT in the
emulsion, followed by spin mixing to dissolve PEI in chloroform phase with a spinning rate of
50rpm. After 5 hours, the resultant mixture of CNTs/PEI/emulsion was treated by high power
ultrasonication for 30mins in iced water bath. The film samples with a thickness of ca. 40μm were
prepared via cast coating method.
7.2.4 Examination of surface modification and dispersion analysis
The effect of non-covalent surface modification on dispersion was examined by two approaches:
(1) Film coating and (2) Centrifugation. In film coating approach, the dilute CNT/emulsion solution
prepared according to the procedure in section 7.2.3 was coated on a glass substrate. For the surface
modified CNT, because the surfactant could effectively change the surface properties of chemically
inert CNT surface and prevent aggregation, a homogenous thin CNT thin film was expected after
surface modification. While the centrifugation could be used to testify the strength of interaction
between surface modified CNT and water, and consequently the stability of surface modification,
both of which are critical to dispersion of CNTs in solvents and polymers. The centrifugation
started from 3000rpm for 5mins, followed by 5000rpm and 7000rpm for 5mins, respectively.
Scanning electron microscopy (FEI 200) was also applied to study the microstructures of CNT thin
films as well as the dispersion of surface modified CNT in PEI matrix.
7.2.5 Electrical properties
Both volume conductivity and surface conductivity were studied. The volume conductivity of
PEI/CNT nanocomposites were measured by both Keithley 6517A with 8009 test chamber and
Keithley 2410 source meter with clip-on test leads for measurement of volume conductivity along
thickness direction and length direction, respectively, while the surface conductivity was measured
by Keithley 6517A with 8009 test chamber. In particular, the volume electrical property along
length direction was conducted on a 50mm (Length) × 20 mm (Width) × 0.04 mm (Thickness)
sample, and two ends of the sample was coated with conductive glue to ensure the contact with test
leads as well as minimize contact resistance.
7.3 Results and discussion
7.3.1 Effect of non-covalent dispersion on CNT dispersion
Figure 7.1 CNT films coated on glass substrates
The surfactant PEDOT: PSS could form non-covalent bond with CNT via π-π stacking, while its
water solubility can lead to the interaction between CNT and water and improve dispersion and
solubility of CNTs in aqueous solution. In order to verify this, we coated aqueous solutions of P-
CNT and S-CNT on glass substrates, as shown in Fig.7.1. Despite the exactly same
ultrasound treatment applied to both CNT aqueous solutions, the differences of the morphologies of
two CNT films are distinct, due to the non-covalent surface modification. In P-CNT film, after the
P-CNTs still tend to aggregate due to high surface energy of P-CNTs, thus there is no continuous
CNT film formed on the glass substrate; on the other hand, the non-covalent surface modification
exhibited remarkable effect on dispersion of CNTs. S-CNTs could form a continuous and uniform
film on the glass, and there are no CNT agglomerates like the ones in P-CNT observed, suggesting
the successful surface modification of CNT via PEDOT:PSS. Furthermore, SEM was applied to
investigate the microstructures of two CNT films.
S-CNT film (Cont.)
Figure 7.2 Morphology of CNT films observed via SEM
According to Fig.2, there are a lot of large P-CNT agglomerates with irregular shapes on the
substrate. After drying, the ultrasonication treated P-CNTs still strongly entangled. The S-CNT film
has completely different structures. Due to existence of PEDOT: PSS, the S-CNT can readily spread
over the substrate to form a uniform S-CNT film. Within the S-CNT film, the CNT was coated and
interconnected via PEDOT:PSS. At the same time, we also noticed a lot of spheres with a diameter
on the S-CNT film. The zoom-in SEM image revealed that these spheres were the S-CNT
agglomerates which might form during drying, however, compared with the agglomerates of P-
CNTs, the S-CNT agglomerates have a much smaller size and regular shapes, which should be also
a result of non-covalent surface modification via PEDOT: PSS. In a word, the non-covalent surface
modification of CNT via PEDOT:PSS could successfully improve dispersion of CNTs.
In addition to the investigation of morphology and microstructures of CNT films, centrifugation
was also applied to aqueous suspension of CNTs to analyze the stability of P-CNT and S-CNT in
water, i.e. the interaction between CNTs and water. Fig.7.3 shows the pictures of both P-CNT and
S-CNT aqueous suspensions after centrifugation. Due to the centrifugal effect and higher mass
density of CNTs compared with water, if the interaction between CNTs and water is not strong
enough, the CNTs will gather at the bottom of the centrifugation tube. Obviously, the P-CNTs do
not have attractive interactions due to the hydrophobic nature of SP2 hybridized carbons, and after
centrifugation, almost all P-CNTs were at the bottom. In contrast, S-CNTs still stably suspended in
water, and they were hardly to find at the bottom of the centrifugation tube, indicating the strong
interaction between S-CNTs and water as a result of PEDOT:PSS. This result does not only prove
the effect of non-covalent surface modification on uniform dispersion of CNTs, but also revealed
the strong non-covalent bonding between CNT and PEDOT: PSS. The strong centrifugation could
not break this non-covalent bonding, and consequently, the S-CNT aqueous suspension still stayed
stable and uniform.
Figure 7.3 P-CNT and S-CNT in water after centrifugation
7.3.2 Microstructures of porous PEI/PEDOT:PSS/CNT hybrid composites
Figure 7.4 SEM images of porous structures of PEI/PEDOT:PSS/ 3.0wt% CNT composites and
dispersion of S-CNTs
The hybrid composites were fabricated by emulsion foaming technique with chloroform and
water as oil and water phases, respectively. The high power ultrasonication could break down water
droplets and disperse them in oil phase in the emulsion. According to Fig.7.4, we can see the
diameter of most pores is around 3μm, with maximum pore size of ca. 6 μm, due to the
homogenization effect o ultrasound. At the same time, it is evident that the S-CNTs could
uniformly disperse in PEI matrix with no obvious agglomerates observed. The fabrication of porous
PEI hybrid composites have some advantages, including reduced mass density as well as uniform
dispersion of S-CNTs due to the existence of water phase.
7.3.3 Anisotropic volume electrical properties of PEI hybrid nanocomposites
0 1 2 3 4 5 6 7 8 9
PEI+0.5 wt% P-CNT
PEI+0.5 wt% S-CNT
PEI+3.0 wt% P-CNT
PEI+3.0 wt% S-CNT
Figure 7.5 Effects of concentration and surface modification of CNTs on the volume electrical
properties of PEI nanocomposites along thickness direction
0 1 2 3 4 5 6 7 8 9
Length directionPEI + 3.0 wt% S-CNT
Figure 7.6 Comparison of volume electric properties of PEI/3.0wt% S-CNT nanocomposites along
thickness and length directions
The addition of 0.5 wt% P-CNT in PEI matrix only slightly decreased the volume electrical
resistivity along thickness direction compared with pure PEI, suggesting that the percolation
network of CNTs did not form at this loading level in PEI matrix. The non-covalent surface
modification of CNTs via PEDOT: PSS had some positive effects on improving the volume
electrical properties, resulting in the decrease of volume resistivity by 2 orders along thickness
direction. At 3.0 wt % CNT loading, the volume electric resistivity along thickness had a dramatic
decrease to 107 ohm·cm, indicating the formation of conductive network. However, the effect of
surface modification on the electric property along thickness direction was not obvious at 3.0 wt%
CNT loading, which might be related to the film coating method and will be discussed later.
Regarding the volume electrical properties along length direction, only the results of PEI / 3.0
wt% S-CNT nanocomposites are given in Fig.7.6. This is because the percolation network of CNTs
has not form in 0.5 wt% CNT reinforced PEI nanocomposites, which led to high resistivity; at the
same time, for PEI/0.5wt% P-CNT nanocomposites, its volume resistivity along length direction is
out of the testing range of Keithley 2410 source meter, suggesting its volume resistivity along
length direction is no less than 107 ohm·cm.
In PEI/3.0wt % S-CNT nanocomposites, the anisotropic volume electrical property was
observed. The volume resistivity along length direction was as high as ca. 8000 ohm·cm, which is
nearly 4 orders higher than that along thickness direction, and also much higher than that in
PEI/3.0wt% P-CNT. These results indicate the contribution of non-covalent surface modification
via PEDOT: PSS on the formation of conductive network in PEI matrix along length direction,
which is also the cast-sliding direction during making film samples.
Figure 7.7 Schematic illustration of mechanism of enhanced electrical conductivity in
PEI/PEDOT:PSS/CNT nanocomposites. R1 represents the non-conjugative polymer block in the
The mechanism of this enhanced electrical conductivity can be explained by the schematic
illustration. The PEDOT:PSS with conjugative structures is attached to the CNT surface via π-π
stacking, and the conjugative structures can also transport electrons. Thus, the attachment of
PEDOT:PSS on CNT surface provides extra conduction path for electrons, which could lead to the
improvement of electrical properties of PEI/CNT nanocomposites. In addition, the discussion in
previous sections also revealed the uniform dispersion of S-CNT in PEI matrix, thus, both uniform
dispersion and extra conduction path as a result of non-covalent surface modification could form a
complex conductive network composed of both uniformly dispersed CNTs and PEDOT:PSS
macromolecules, which could efficiently increase the electric conductivity.
However, according to Fig.7.6, the non-covalent surface modification did not show obvious
effect on the volume electrical properties along thickness direction in 3.0 wt% CNTs reinforced PEI
nanocomposites. This may be explained the film coating procedure. In this study, a cast-sliding
method was applied to make the film samples. The sliding process may lead to stretching and
extension of PEDOT: PSS polymer chains along length direction, which help the formation of the
complex conductive network, and did happen in thickness direction. In 0.5 wt% CNT reinforced
PEI nanocomposites, the more remarkable effect of surface modification should be related to the
formation of percolation network. Because percolation network has not formed yet in P-CNT
reinforcement PEI nanocomposite at 0.5 wt% loading, and the PEDOT: PSS attached on CNT
surface could assist the formation of percolation network to some extent and decrease the resistivity
more or less. In 3.0wt% CNT reinforced PEI nanocomposites, the percolation network has been
already constructed, thus the effect of non-stretched PEDOT: PSS might be ignored.
7.3.4 Surface electrical properties
0 1 2 3 4 5 6 7 8 9
PEI+0.5 wt% P-CNT
PEI+0.5 wt% S-CNT
PEI+3.0 wt% P-CNT
PEI+3.0 wt% S-CNT
Figure 7.8 Surface resistivity of PEI/CNT nanocomposites with and without non-covalent surface
The mechanism illustrated in Fig.7.7 can also explain the surface electrical properties of
PEI/CNT nanocomposites. With addition of only 0.5 wt% P-CNT, the surface resistivity only
slightly decreases by 1 order, while the non-covalent surface modification results in 3 more orders
decrease, which makes the surface resistivity of PEI/0.5wt% S-CNT close to PEI/3.0wt% P-CNT
nanocomposites. Due to higher electrical conductivity of CNT, compared with carbon nanofiber
(CNF) , P-CNT could lead to better surface electrical properties of PEI nanocompsoites. Thus,
at 3.0wt% CNT loading, the surface resistivity can reach ca. 1011
ohm·sq, which is slightly lower
than the minimum requirement of antistatic application (1012
ohm·sq). The surface modification via
PEDOT:PSS shows further decrease of surface resistivity, as shown in Fig.7.8. The surface
resistivity of PEI/3.0wt% S-CNT is as low as ca. 5×106
0 1 2 3 4 5 6 7 8 9
PEI+ 4wt% PEDOT:PSS
Figure 7.9 Electrical properties of PEI/PEDOT:PSS blend
As discussed earlier, the formation of complex conductive network by uniformly dispersed
CNTs and PEDOT:PSS conductive paths resulted in the dramatic electrical properties of PEI/CNT
nanocomposites. In order to verify this mechanism, we also testified the electrical properties of
PEI/PEDOT:PSS blend with high PEDOT:PSS concentration. Because the PEDOT:PSS used in this
study have a maximum electrical conductivity of 1S/cm according to the provider. Thus, it is also
reasonable to speculate the increased conductivity of PEI/S-CNT nanocomposites also comes from
high conductivity of PEDOT:PSS. However, according to Fig. 7.9, with 4wt% PEDOT:PSS in PEI
matrix, both surface resistivity and volume resistivity are much lower than those of PEI/3.0wt% S-
CNT, in which the concentration of PEDOT:PSS is 3.75% according to the ratio of CNT to
PEDOT:PSS. This result further proves the contribution of the complex conductive network to the
high electrical conductivity in PEI/S-CNT nanocomposites.
The non-convalent surface modification via poly(3,4-ethylenedioxythiophene)-
poly(styrenesulfonate) (PEDOT:PSS) was successfully applied to CNT and CNT reinforced PEI
nanocomposites. It proves to be an efficient way to improve the dispersion quality of CNT with
high surface energy and electrical properties of PEI / CNT nanocomposites. With addition of only
3.0wt% surface modified CNT, the volume resistivity and surface resistivity of PEI nanocomposites
can reach really low levels: ca. 8000 ohm·cm and 5×106 ohm·sq, respectively, due to the enhanced
conductive network in PEI matrix. This research indicate the great potential of non-covalent surface
modification in fabricating advanced polymer nanococmposites with superior functionalities, which
could broaden applications of polymer nanocomposites in airplane, ground transportation,
electromagnetic shielding, construction as well as electronics that requiring high conductivity and
static dissipation capability. Furthermore, the non-covalent surface modification is less complicated
than traditional covalent surface modification, and fewer toxic and volatile chemicals are involved ,
which satisfy the demands of green manufacturing technique in polymer industry.
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Conclusions and Future Plans
In this study, the modification of polyetherimide (PEI) by graphitic carbon nanofillers (GCNs)
has shown remarkable success in both improvement of fundamental properties and new
functionalities. A series of significant research findings have been revealed:
1. A negative permittivity phenomenon in a class of PEI/carbon nanofiber composite films with
strong composition dependence was discovered. This finding will boost the development of flexible
and lightweight metamaterials which may be used to render matter and antennas invisible.
2. A bi-layer nanocomposite with more effective static dissipation performance at a very low
loading level (0.6wt%) was designed. This bi-layer nanocomposite exhibited a static dissipation rate
1000 times higher than conventional monolayer nanocomposites, while the nanofiller loading was
reduced by 91.4wt%, significantly benefiting aero applications that require high antistatic capability
and fuel efficiency.
3. Effective dispersion of carbon nanofiber in PEI matrix could dramatically improve
mechanical and tribological performance: strength and toughness increased by 50% and 500%,
respectively, and the wear rate decreased by 56%, while using a free-chemical and industry friendly
4. Covalent surface modification was applied to modified graphene nanoplatelets via
organosilane treatment, and the resultant PEI nanocomposites possessed high damping capacity.
The damping factor was 3 times higher than that of the pure polymer, while the modulus increased
4 and 200 times at 30°C and 200°C, respectively, suggesting the potential of graphene nanoplatelet
composites for acoustic damping applications.
5. Non-covalent surface modification of carbon nanotubes via conductive polymers could
dramatically increase both surface and volume conductivities of PEI nanocomposites at a rather low
loading level by 5 orders and 4 orders of magnitude, respectively, compared with pristine carbon
nanotube reinforcement nanocomposites.
6. A novel but facile non-destructive approach for quantitatively evaluating nanofiller dispersion
in a polymer matrix, which is applicable in industry for quality control of nanocomposite products
via AC/DC conductivity ratio, was also developed.
However, while we are celebrating these achievements, the existing problems are also obvious.
First of all, although some properties, such as damping properties, have showed remarkable
improvement by GCN reinforcement, they still do not meet the requirements for practical
applications; secondly, the application of surfactants in polymer nanocomposites showed many
benefits in improving properties and functionalties. However, they might also cause degradation of
certain properties; for example, PEI/GCN nanocomposites modified by organosilane and by
conductive polymers with low thermal stability will have reduced thermal stability compared with
pristine GCN reinforced polymer nanocomposites.
Thus, solving these existing problems will be the main tasks of future work. In order to
substantially improve the properties and functionalties of PEI/GCN nanocomposites, the following
efforts will be made:
1. Bi-layered nanocomposite structures, hybrid GCN reinforcement, and non-covalent surface
modification will be optimized and integrated to develop new fabrication techniques for
nanocomposites. Each technique has been effective in improving properties and decreasing
GCN loading of PEI/GCN nanocomposites. Thus, the optimization of each technique will be
carried out, in order to determine the ideal nanocomposite processing parameters. It is also
believed that the combination of any two or all three of these techniques will improve the
properties of PEI/GCN nanocomposites to new levels and make them suitable for practical
2. New surfactants will be developed. Proper surface modification is necessary in many
applications, due to its effect on some functionalties. In particular, conductive polymers
have exhibited promising applications in functional polymer nanocomposites. However,
their low degradation temperature and brittleness could degrade the thermal and mechanical
properties of polymer nanocomposites. Thus, new conductive copolymers with higher
thermal and mechanical stabilities will be developed in order to maintain the fundamental
properties of PEI/GCN nanocomposites.
Finally, the pursuit of environmentally friendly nanocomposite manufacturing techniques is the
central issue of this research project. Therefore, we will try to use less toxic and less volatile
chemicals, and reduce the energy consumption during manufacturing by controlling the processing
temperature and pressure. Instead of he complicated physical and chemical procedures commonly
used in nanocomposites processing, simpler processing techniques will be developed without
sacrificing the properties and functionalities of PEI/GCN nanocomposites.
Other Research projects in addition to dissertation project
Novel Hydration Induced Flexible Sulfonated Poly(etherketoneketone) Foam with
Super Dielectric Characteristics
Journal of Materials Chemistry, 2011, 21, 13546.
Highlighted by Health and Medicine Week, September, 19th
, 2011, Page 1792.
“Findings from Washington State University Broaden Understanding of Materials Chemistry”
High performance thermoplastic polymers, such as poly(etherketoneketone) (PEKK) etc., with
excellent mechanical properties and thermo-oxidative stability are in great demand for expanding
commercial applications. Transforming these materials into foamed structures through energy
efficient approaches can be dramatically significant for various applications in which weight is
critical. Due to the rigid aromatic structures, foaming them is normally extremely difficult via
blowing agent assisted foaming techniques. In this study, homogeneous sulfonated PEKK (SPEKK)
foam was fabricated through a facile and energy efficient hydration reaction induced self-foaming
approach, in which, water acted as the blowing agent. The resulting SPEKK foam has a uniform
pore size of ca. 5µm and cell denisty of 9.77×109cell/cm
3, as well as a very low mass density of
0.42g/cm3, suggesting a superiority to other foam structures reported, such as poly(aryl ether ketone)
microceullar foams fabricated by conventional blowing agent assisted foaming techniques. The
dielectric study demonstrated that the microcellular SPEKK foam possessesd great potential for use
in energy storage applications. The as prepared SPEKK foam showed a dramatically high
permittivity (109 at 0.01 Hz and ca.10
6 at 1000 Hz). At the same time, the SPEKK foam has a rather
low dielectric loss, between 2~10 over the whole frequency range. With the dehydration of SPEKK
foam, the relative permittivity dropped, but was still as high as ca.107 at 0.01 Hz and ca.10
High performance thermoplastic polymers, such as polyimide, polysulfone poly(aryl ether
ketone) (PAEK) etc., with excellent mechanical properties and thermo-oxidative stability are in
great demand in various applications, in particular for flexible, low mass density and high
temperature energy applications. Thus, energy materials based on these polymers have been of great
interest today [1-9]. The family of PAEK includes a series of polymers consisting of aromatic ether
groups and ketone groups with varying ketone/ether (K/E) ratios. As a result of plentiful phenyl
groups on polymer backbones, they possess high modulus and strength as well as excellent thermal
stability. The research on potential applications of PAEK falls into two main categories. Firstly, the
PAEKs and their composites are commonly considered as ideal biomedical materials [10-12];
secondly, the PAEKs and sulfonated PAEKs (SPAEK) have attracted great interest in energy
applications, including the energy storage applications of PAEK , and proton exchange
membranes of SPAEK for fuel cells [1,3,4,6-9]. Polyetherketoneketone (PEKK) with a K/E ratio of
2:1, is gaining more and more attention in the application of proton exchange membranes, because
the high K/E ratio benefits the oxidative stability and allows for high sulfonation levels . The
SPAEKs including SPEKK are actually polymer electrolytes which could help the transport of
protons (hydrogen ions) and consequently exhibit proton conduction [1-9]. Owing to this property,
the proton conduction of SPAEKs has been the research focus for many years. The effect of water
intake [9,14], sulfonation degree [1,6,9,14], temperature [3,9,14], acid doping  as well as the
molecular structures[9,14] of SPAEK on the proton conductivity have been extensively studied,
whereas other potential applications/properties of SPAEK in the energy field have been rarely
explored, for instance, energy storage/dielectric properties.
Porous polymers (polymer foams) represent a group of functional materials with extremely low
mass density and high surface area. They are becoming more important in energy applications, due
to extended performances, deformability as well as the new opportunities depending on
structural/functional properties .The current fabrication techniques of polymer foams basically
include three types. The mostly classic one is using a blowing agent assisted foaming process, in
which, the gas generated by decomposition of blowing agent, such as Al(OH)3 , or thermal
expansion of saturated inert gas, such as N2, Ar, CO2[17-20] that could expand the free volumes in
the polymers and lead to the formation of porous structures. The second one is template-assisted
approach, in which the multiphase polymer composites/blends are fabricated, and the porous
structure will form after removing the dispersed phase [21,22]. The third one is to synthesize
polymers with large free volume, such as hypercrosslinked polystyrene  and spirobifluorene
derived polyimide with intrinsic porosity .
For high performance aromatic polymers, such as polyimide, polysulfone as well as PEAKs, etc,
their foams have been highly demanded in new generation electronics, aerospace, transportation
industries as well as energy applications, due to their low mass density, low dielectric constant,
desirable mechanical and thermal properties . However, their extremely high melt viscosity, glass
transition temperature (Tg) and melting point, make the foaming process more challenging.
Currently, CO2 is most frequently used as the blowing agent for these high performance polymers
[17-20]. However, this foaming procedure requires high temperature, high pressure as well as
complicated processing controls. Specifically, this foaming technique includes three steps: first, the
preparation of solid core polymer and polymer composites samples; second, saturation of polymers
with CO2 at elevated pressure (e.g. 810-820 psi for polyethersulfone and polyphenylsufone ) for
sufficient long time (a few hours to several days); third, foaming of saturated polymers at
high temperature, usually around glass transition temperature. In these procedures, there are many
parameters, including saturation time, saturation temperature, foaming time and foaming
temperature, etc., that make the quality control of the polymer foams more complicated and difficult.
In particular, for large foamed parts, the control of porous structures from its surface to the core is
more challenging. In fact, though all three of the foaming techniques mentioned have shown
success in fabricating foams of these high performance polymers [15-24], high energy consumption
and low efficiency are always needed, and the fabrication procedures, especially, the template-
assisted approach and synthesis of polymers with large free volume, are usually accompanied with
toxic chemicals [21-24].
Due to the rigid polymer structures of the PAEK series, foaming them is extremely difficult via
blowing agent assisted foaming techniques. Thus far, foaming of PAEKs has not been widely
studied as have other high performance polymers with high Tg and melt viscosity, for instance,
polyimide [21, 25, 26] and polysulfones [20, 27]. Through conventional approaches of applying
blowing agents, only large pore sizes and wide size distribution that resulted in less reduced density
(closed to the solid polymer level) have been obtained for PEEK. For example, with blowing
saturated inert gas, the PEKK foam with the mass density of 0.52g/cm3
~ 0.65g/cm3 has been
fabricated, and its pore size was between 10 ~100μm. R. Verdejo et al  used two blowing
agents, Al(OH)3 and Clariant XH907, to prepare injection molded foams of PEEK and
PEEK/carbon nanofiber composites. For the pure PEEK foam, both blowing agent led to a large
pore size and wide size distribution in the range of 60 ~700 μm , while the mass density of the
foams was pretty close to the solid PEEK: ~ 1.15 g/cm3 for foams vs. ~1.25 g/cm
3 for solid
PEEK. A study on the nanocellular foam of PEEK/PEI blends via blowing agent CO2 revealed
that under the same foaming conditions, only very few pores appeared in PEEK, while PEI
exhibited homogeneous nano porous structures . Therefore, the successful fabrication of PAEK
and SPAEK foams with well-controlled pore size and low mass density are still scant. Also, a facile,
energy efficient as well as less toxic foaming approach is necessitated for this group of high
In this study, we proposed a new self-foaming approach via a hydration reaction to achieve high
performance flexible SPEKK microfoams. The whole foaming process was accomplished at room
temperature and one atmospheric pressure within a few minutes, representing a new efficient and
environmentally friendly technique to fabricate porous polymer structures. The only chemicals
involved in this self-foaming procedure are concentrated sulfuric acid and water, while the only
byproduct is the recyclable dilute sulfuric acid. The SPEKK foam obtained via the approach used in
this study has much more homogeneous and smaller porous structures (pore size around 5μm), as
well as rather low mass density of 0.42g/cm3. Compared to the reported studies on high temperature
polymer foams [17-27] fabricated by various foaming techniques , the resulting SPEKK foam
prepared through the facile approach can be promising for signifciantly extending the application of
these materials. In particular, their dramatic high relative permittivity and extremely low loss reveal
the great potential of SPEKK foams in high efficiency energy storage applications.
2.1.Materials and preparation of sulfonated PEKK microfoam
8 grams poly(etherketoneketone) (PEKK) particles provided by Cytec Industries Inc. were
mixed with 50 ml concentrated sulfuric acid. As a result of sulfonation process, the PEKK could
dissolve in concentrated sulfuric acid and form sulfonated PEKK (SPEKK). According to the
previous works on the sulfonation kinetics in mixture of concentrated and fuming sulfuric acid ,
the sulfonation degree will stay relatively stable after 60 hrs at 40°C, thus, in order to obtain a high
sulfonation degree in this study, the SPEKK/sulfuric acid solution was sealed at room temperature
for 5 days to assure the sulfonation reactions. The resulting SPEKK solution has a “syrup” look, as
shown in Fig. 1(A). The preparation of SPEKK microfoam was accomplished by the procedures
illustrated in Fig.1 (B) – (E). A glass rod was dipped into the viscous solution of SPEKK/sulfuric
acid (Fig.1(B)), and then transferred to water (Fig.1(C)). Because of the poor solubility of SPEKK
in water and diluted sulfuric acid, the SPEKK will precipitate and form a milk-white thin sheet with
thickness of 0.6 mm-1.0mm wrapping around the glass rod (Fig.1(E)), which is found to be the
SPEKK microfoam by microstructure analysis. The foaming mechanisms will be introduced in
details in following section (Fig.1(E))
Water absorption is a typical property of SPEKK due to the existence of hydrophilic sulfonic
groups (-SO3H-), with the aggregation of which forms the domains acting as water reservoirs .
The water absorption strongly depends on the air humidity, temperature and degree of sulfonation,
and significantly affects the mechanical and physical properties of SPEAKs [6, 28]. Therefore we
investigated the effect of the water content on the conductivity and dielectric properties. Three
samples with different water contents have been studied, and the results are: (1) As-prepared
SPEKK foam (~77.8wt% water) did not experience any drying treatment before test, and only the
sample surface was dried for the convenience of tests. The test started within 10 min. since SPEKK
was taken out of the water. It should be pointed out because of the porous structure of SPEKK foam,
the ~77.8% water includes both free water in the pores and absorbed water in the “water reservoirs”
in SPEKK matrix. (2) Dry -1 SPEKK foam (~ 14.6% water) was obtained by exposing wet sample
to the air with a relative humidity of approximate 30% till the water content stabilized at a constant
level. (3) Dry 2 SPEKK foam (~2.0wt% water) was obtained by oven drying as-prepared SPEKK
foam at 70°C until the sample weight did not vary (up to 2 hrs). Although the completely dried
sample has no water inside, the highly hydrophilic SPEKK foam will absorb water quickly after
taken out of the oven. The 2wt% water content of this sample was measured right after conductivity
and dielectric analysis was done, which guaranteed extremely low water content during the test. The
water content was calculated by equation (1) :
where Wwet is the weight of the samples containing water and Wdry is the completely dried
sample. The SPEKK foam in this study has a moderate water absorption level, compared with other
SPAEKs . According to reported work, the water content could range from only below 10% to as
high as over 100% [1,6,9], while mostly falling into 20%~70%.
2.2.Characterizations and tests
The foam structures were studied by Scanning Electron Microscope (SEM) with image analysis
software. Thermal gravity analyzer (TGA) was applied to investigate the mass loss of the SPEKK
foam with temperature at a heating rate of 20°C/min. From the thermal decomposition, the
sulfonation degree of SPEKK can be evaluated. Differential scanning calorimetry (DSC) at a
heating/cooling rate of 10°C/min was applied to obtain the glass transition temperature (Tg). The
proton conductivity, relative permittivity and dielectric loss were analyzed by Universal Dielectric
and Impedance Spectrometer BDS 20 from Novocontrol Inc. The circular samples with a diameter
Figure 1 (A) The syrup like SPEKK/H2SO4 solution; (B)-(E) Procedures
and mechanism of preparation of SPEKK foam
of 20mm were cut from the bulk SEPKK foams using a sharp blade. Before testing, the sample
surface were carefully cleaned with acetone, and then coated by carbon based conductive glue to
improve the effective contact as well as minimize the interfacial polarization between electrodes
and sample surfaces. The proton conductivity, dielectric properties ranging from 0.01 Hz to 106Hz
have been investigated at room temperature.
3. Results and Discussion
3.1. Hydration reaction induced self- foaming process, structures, and mechanisms
Right after immersing the SPEKK/H2SO4 viscous solution into to the water, the outermost layer
turned to white solid immediately, due to the precipitation of SPEKK in water and diluted acid. The
SPEKK and/or other SPEAKs are usually insoluble or just swollen in water and diluted H2SO4
solution, although the sulfonic groups on SPEKK chains are highly hydrophilic. As illustrated in
Fig.1 (D), the concentration gradient of sulfuric acid exits between the syrup-like SPEKK solution
and water (Fig1.(D)), thus, the diffusion of H2SO4 happens as a result of the concentration gradient,
as does H2O. The resulting lowered concentration of H2SO4 in SPEKK solution forced the SPEKK
out of the liquid phase and precipitate to milk-white solid (Fig.1 (E)). The diffusion of H2SO4 and
H2O proceeded very fast. Within 10 min, the 0.6mm ~ 1.0mm thick thin sheet of SPEKK was
obtained. The SEM morphology studies revealed that the white solid had a rather uniform
porous/foam structure which possessed superior flexibility, as shown in Fig.2. According to
statistical analysis, some critical parameters describing foam structure are given here: cell diameter
of ~ 5µm; cell density of 9.77× 109cell/cm
3. The mass density of completed dried SPEKK foam is
only ca. 0.42g/cm3.
Obviously, these foaming results indicate that the SPEKK foam obtained in this study led to
remarkably homogeneous structures with small pore size and more uniform size distribution, as
well as low mass density (the relative density to the solid polymer is only 0.36), as compared with
the reported studies, which show very low porosity in PEEK , broad pore size 10 ~100μm ,
a large pore size and wide size distribution 60 ~700μm while a mass density close to the solid
PEEK . All these reported forming results were via conventional blowing agent techniques.
Even compared with other polymer foams, such as microcellular and nanocellularPEI foams ,
microcellular polysulfone foams [20, 27] and microcellular polystyrene foam [17, 30], the mass
density of the SPEKK in this study is still at a rather low level.
In order to understand the foaming mechanism, we can look back into the three types of
foaming approaches previously mentioned, and it is easy to exclude the possibilities of template
Figure 2 (A) Flexible SPEKK foam and (B) EM images of porous
structures of SPEKK foam with different magnifications
assisted foaming and synthesis of polymer with large free volume, since there were no multiple
phase structures and polymerization procedures introduced during the preparation of SPEKK foams.
If the foaming of SPEKK falls into the category of blowing agent assisted foaming technique, it is
necessary to find out the blowing agent and driving force to blow the polymer matrix.
In this foaming process, only PEKK particles concentrated H2SO4 and diluted H2SO4 solution
were involved. The same as precipitation of SPEKK, the foaming of SPEKK is also believed as a
result of the diffusion of H2SO4 and H2O. The diffusion does not only lead to the changes in the
concentration of H2SO4, but also induces the thermodynamically favorable hydration reactions of
H2SO4, as shown in Fig. 1(D). The hydration reactions of H2SO4 are highly exothermal. When the
concentrated H2SO4 is added to water, on one hand, the water will be protonized. On the other hand,
an enormous amount of heat will be released, accompanied by a dramatic increase of temperature of
the H2SO4 solution. It was found in this study that the liquid temperature increased up to 110°C,
when 50ml concentrated H2SO4 was added to 400 ml water at room temperature. The fact that this
temperature is above the boiling point of water at one atmospheric pressure suggests the rapid
evaporation of water. When the H2O molecules diffused into the SPEKK/ H2SO4 solution, it can be
seen as the addition of water into concentrated H2SO4, accompanied with the occurrence of
hydration reactions and resulted in a large amount of heat able to vaporize the water. In the
SPEKK/H2SO4 solution, a lot of free volume exists among SPEKK polymer chains, at the same
time, compared with polymer melt, the viscosity of the solution is much low, therefore, the rapid
evaporation and expansion of water vapor can easily expand these free volumes to form the porous
structures, while the simultaneous precipitation of SPEKK as a result of lowered H2SO4
concentration solidify these porous structures and lead to the formation of SPEKK foam. Thus,
vapor water acts as the blowing agent, and the heat generated via hydration reaction of H2SO4
provides the energy to vaporize water and thermally expand the free volume. It is also speculated
that the viscosity of the SPEKK/H2SO4 solution plays an important role in the foaming procedure
by affecting the thermal expansion of free volume in viscous liquid, which will be further studied.
The significant role of water in this foaming process can be further verified by replacing water with
dichloromethane (CH2Cl2) as shown in Fig.3. The foaming of SPEKK did not occur in CH2Cl2
solvent, due to 1) CH2Cl2 is immiscible with concentrated H2SO4, thus there is not the same
diffusion phenomenon as illustrated in Fig.1, and 2) the SPEKK still stays in concentrated H2SO4
region as the syrup-like viscous solution, instead of precipitating in CH2Cl2, at the same time, the
hydration reaction does not exist between CH2Cl2 and concentrated H2SO4. Consequently, there is
no heat source for thermal expansion of free volume of polymers. In summary, the foaming of
Figure3 Evolution of viscous SPEKK/H2SO4 solution in water and Dichloromethane (CH2Cl2)
SPEKK can be described as a self-foaming process induced by hydration of concentrated H2SO4,
since the concentrated H2SO4 has also been used to sulfonate the PEKK to obtain SPEKK.
The advantages of this foaming procedure are obvious, including (1) high energy and cost
efficiency, (2) environment friendliness (3) facile operation. The whole process is accomplished at
room temperature and one atmospheric pressure, and correspondingly the expensive and
complicated temperature and pressure control systems are not needed. There are no toxic and/or
costly monomers, chemicals and blowing agents involved, as well as no complicated chemical
reactions taking place; at the same time, the foaming process does not yield any toxic byproducts.
The only byproduct is diluted H2SO4 aqueous solution which can be easily collected and recycled
for next foaming process or other applications. Compared with other foaming approaches, it has less
risk to the environment and operators. And lastly, since the aromatic sulfonation commonly exists,
this approach shows great potential in the preparation of aromatic thermoplastic foams.
0 100 200 300 400 500 600 700 800 9000.3
Figure 4 TGA curves of sample dry-1 SPEKK foam
3.2.Thermal Analysis of SPEKK foams
The TGA was applied to analyze the thermal properties of the SPEKK. Fig.4 shows the weight
loss curve of dry -1 SPEKK foam (~14.6wt%) vs. temperature at a heating rate of 20ºC/min.
Obviously, the weight loss was related to water loss at temperatures below 100ºC, similar to the
PVA/PAA membrane containing free water , disulfonation of sulfonic groups around 350ºC, as
well as the decomposition of backbone above 500 ºC . At the same time, the DSC results
revealed the increase of glass transition temperature (Tg) from 158oC of PEKK to 163°C of SPEKK,
is a result of sulfonation. It was found that the higher degree of sulfonation could lead to larger
increase in Tg, and an increase over 10 o
C has been reported [1, 6, 28]. The 5 °C increase in our
study may suggest that the further sulfonation is still needed, via longer reaction time and higher
Protons (hydrogen ions), as the only free charge carriers in the SPEKK foams studied here,
account for the AC conductivity as shown in Fig.5. For SPAEKs, their proton conductivity strongly
depends on water intake, sulfonation degree and temperature . Usually, the high water content,
high sulfonation degree as well as high temperature could effectively lead to the high proton
conductivity. In this study, the highest conductivity of ca. 5.0E-4 S/cm was observed in a wet
sample with 77.8% water, which is 1-2 orders lower than those reported solid SPAEKs [1,4,6,9,].
A study on the SPEEK with 72% sulfonation degree and 52% water showed a proton conductivity
of 2.8E-3S/cm at room temperature , and another similar study on SPEEK with 65 sulfonation
degree and 20% water showed a proton conductivity of 5E-3S/cm. Other than the differences in
testing apparatus/procedures and molecular structures of PEKKof this study , it is believed that
reasons for lower conductivity of porous SPEKK are the moderate sulfonation degree and porous
structure. First of all, the DSC results already revealed that the sulfonation degree of SPEKK foam
in this study was not high enough, according to the fact of 5 oC increase in Tg, as discussed in
previous section, which might be improved by increasing the sulfonation temperature . On the
other hand, owing to the porous structure in the wet sample, the 77.8% water must include a portion
of free water in the pores, which has no significant contribution to the proton conduction. As a
result of dehydration of two dry samples, the conductivity drops to 2E-6S/cm for dry-1 SPEKK
foam (14.6% water) and 1E-5 S/cm for dry-2 SPEKK foam (2% water), as expected. The lower
conductivity of dry-1 SPEKK foam (14.6%water) compared with dry-2 SPEKK foam (2% water)
might be due to the diffusional loss of protons when exposed to the air. On the other hand, during
oven drying of the as-received sample at 70 o
C, some residual sulfuric acid molecules will further
the sulfonation reactions of SPEKK, which may lead to slightly higher sulfonation degree of dry-2
SPEKK foam (2wt% water)
Figure 5 AC conductivity of microcellular SPEKK foams
3.4.Electrical energy storage /Dielectric properties evaluation
The energy storage capability of materials can be evaluated by electric energy density, which is
directly related to relative permittivity, as given by equations (2). Where, U is energy density
(J/cm3), E is electric field, ε0 is the permittivity of free space and ε’ is the relative permittivity of
materials. The complex permittivity (ε*) and dissipation factor (tan(δ)) are given in equations (3)
1( ' )
2U E (2)
* ' "i (3)
Usually, pure polymers have low relative permittivity in the range of 2-10, which can be
improved by synthesis of copolymers  or fabrication of polymer blends and polymer composites
[31,36-43]. However, with moderate increase of permittivity, the dielectric loss (ε”) exhibits an
unwanted rise, and consequently, increases the dissipation factor, as a result of enhanced
reorientation of polar groups and interfacial polarizations in polymers and polymer composites.
High energy loss does not only reduce the efficiency, the resulting enormous heat also impacts other
properties of the materials and shortens their service life. Till now, the high loss of the energy
storage materials still bothers the scientists working on the high energy efficiency and low loss
materials for energy applications.
Impressive dielectric properties have been found in the SPEKK foams in this study. According
to Fig.6, all the samples show extremely high relative permittivity in the whole testing range. In
particular, the as-received SPEKK foam, which has the highest water intake, has the highest relative
permittivity. At 1 KHz, the relative permittivity is above 105, while at 0.01 Hz, it is even higher
than 109. Both dry samples with differing water content (~14.6wt% and ~2wt%) show moderate
drops in relative permittivity to ca.103 at 1KHz and ca. 10
7 at 0.01 Hz respectively. Obviously,
the extremely high permittivity primarily related to the water. Many researchers have reported that
the water absorption could increase the permittivity of polymeric systems, such as a series of epoxy
composites [44, 45], unsaturated isophthalic polyester/natural fibers composites , as well as the
sulfonated poly (styrene–ethylene/butylenes–styrene) triblock copolymer  etc., because of the
high permittivity of water itself and the interaction between highly polar water molecules with other
groups. In the study of the sulfonated poly(styrene–ethylene/butylenes–styrene) triblock
copolymer, with the increasing of the relative water content to 25wt% H2O/Film, the
> 109, at 0.01 Hz
> 105, at 1KHz
Figure 6 Relative permittivity of SPEKK foams
permittivity went up to 105 at 100 Hz from 2 at 0 wt% water, as a result of the interaction between
the sulfonic groups and polar water dipoles. In this study, the as-received SPEKK foam has a water
content of ca. 77.8%., which is a moderate water content among the SPAEKs, compared with the
reported foamed PEAK series with up to 100% water contents [6, 9]. This amount of water is
considered the main reason for the dramatically high relative permittivity of as-prepared SPEKK
foam. We took advantage of this characteristic of SPEKK to explore the polymeric energy storage
materials used in a water environment with high moisture content which possess practical
applications, but which has not previously received adequate attention.
However, in our SPEKK foams, a great amount of hydrogen ions (H+) also exits in the SPEKK
matrix as a result of simultaneous foaming and precipitation processes, which can also be
considered as a self-doping process. Their contributions can be revealed via the study on the
dielectric properties of dry 2 SPEKK foam with the lowest water content (~2 wt%). The low water
content was reached by oven drying at 70 0C to remove all the moisture within the SPEKK matrix.
Until the end of dielectric test, the water content did not exceed 2wt%. According to Fig.6, on one
hand, the permittivity decreases as a result of the loss of water; on the other hand, the permittivity of
dry -2 SPEKK foam (2 wt% water) is still sufficiently high, and even slightly higher than that of
dry-1 SPEKK foam (14.6wt%water), suggesting the contribution of hydrogen ions. The H+
actually just a proton with really low volume, and can freely move within the SPEKK matrix and
interact with sulfonic groups (-SO3H-) owing to opposite charges. The interaction between H
positive charge and -SO3H- with negative charge could lead to the formation of a great amount of
ion-ion dipoles. Thus, although a smaller amount of water was in the “dry” samples, their
permittivity is still remarkably high. Furthermore, the higher water content could also lead to much
better mobility of H+ -SO3H
- dipoles under an electric field, corresponding to extremely high
Dielectric loss is another important aspect of evaluation of energy applications. The materials
with high permittivity and low loss have always been of great interest. When the orientation and
arrangement of dipoles in the materials result in the high permittivity, they also dissipate a great
amount of energy. Thus, the increase of permittivity inevitably increases the dielectric loss.
Currently, the high loss is still the biggest obstacle to the development of highly energy efficient
materials. Recently, several approaches have been developed to restrain the dielectric loss, such as
introducing an insulating layer between epoxy and silver nanoparticles , three-phase PVDF/
BaTiO3-MWCNT composites  and fabrication of multilayer poly(vinyldiene fluoride)
nanocomposites. In these composites, although the dielectric loss was limited at a very low
level, their permittivity did not show significant improvement as did the SPEKK foam in this study,
being just on the order of several hundred. Fig.7 shows the loss factor (tan(δ)) of SPEKK foam
Figure 7 Dissipation factor of SPEKK foams
samples. It is intriguing that the as-prepared SPEKK foam with the highest permittivity (109 at
0.01Hz) also has the lowest loss, only 2~10 in the broad testing range, indicating that the H2O does
not only account for the high permittivity as discussed earlier, but also successfully inhibits the
dielectric loss. The comparison between the as-prepared SPEKK foam with other reported
polymeric materials with improved dielectric properties was summarized in Table1. Compared with
other polymeric material systems, this advantage of as-prepared SPEKK foam is more outstanding.
Usually, in order to realize the high relative permittivity of polymers, multiple phase
composites/blends are fabricated, as given in Table1 and Fig.8(A). Basically, conductive fillers and
high permittivity ceramic particles are preferable additives for this purpose. However, in these
material systems, the dramatic relative permittivity as high as that of the as-prepared SPEKK foam
has rarely been reported, in particular, polymer foams with such a high relatively high permittivity
have never been reported before, according to authors’ best knowledge; on the other hand, with
moderate improvement in relative permittivity, the dissipation factor (tan(δ)) often has a more
remarkable increase. According to Table 1, for example, the Epoxy/Graphite nanosheet
nanocomposites with a relative permittivity of 250 showed a tan(δ) of 8, while in
LMWPE/UMWPE/carbon nanofiber composites with a relative permittivity of ca. 5000, the
dissipation factor was as high as 1000. The loss of both composites is obviously higher than that of
as-prepared SPEKK foam. For a clearer understanding of the advantage of as-prepared SPEKK
foam, the parameter, ε’/tan(δ), was introduced, which was used to normalize the relative
permittivity ε’ of materials at tan(δ) of 1, and it represents how high the relative permittivity of
materials can be at the same loss dissipation level. Basically, the higher ε’/tan(δ) value indicates
Table 1 Comparison of dielectric properties of as-prepared SPEKK foam with literatures
better efficiency: higher relative permittivity with lower loss, which is ideal for energy applications.
As shown in Fig.8, the as-prepared SPEKK foam shows the highest relative permittivity as well as
ε’/tan(δ) values at various frequencies, suggesting its superior dielectric efficiency to other
composites systems, in particular, to polymer/conductive fillers composites. Compared with other
polyelectrolyte systems, such as (PVA)0.7(KI)0.3·xH2SO4 and [(100-x)PEO + xNH4SCN]:Al-Zn
composites, and PVA/LiFeO4, etc, the as-prepared SPEKK foam in our study still possesses
more impressive dielectric properties, correspondingly, more ideal for energy applications. In
addition, compared with the solid polymeric dielectric obtained via adding rigid and heavy fillers in
polymer matrix, such as those in table 1, the SPEKK foam in this study exhibited superior dielectric
performances without compromising low mass density and flexibility of the polymers.
Log (Frequency) (Hz) ε' Tan(δ) Ref.
As-prepared SPEKK foam (77.8wt% water content)
0 108 2.6
2 2×106 2.9
3 2×105 4.4
4 105 9.2
LMWPE/UMWPE/Carbon fiber 0 ~5×103 ~1000 
PANI/CaCu3Ti4O12 2 106 5-10 
(PVA)0.7(KI)0.3·5H2SO4 2 ~ 4×104 ~ 6 
[93PEO + 7NH4SCN]: 2wt%Al-Zn 2 ~3×105 ~0.5 
Polystyrene/ foliated graphite 2 ~104 ~10
EPOXY/Graphite Nanosheet 3 240 8  PVDF/Graphite Nanoplatelet 3 ~ 5×10
3 ~ 1×10
PANI/EPOXY 4 4×103 0.5 
Obviously, the as-prepared SPEKK foam realized both extremely high permittivity and rather
low dielectric loss. Also, we can observe a relaxation peak between 105-10
6 Hz, which supposes to
be related with the relaxation of hydrogen ion-sulfonic group dipoles. The similar relaxation was
observed in (PVA)0.7(KI)0.3·xH2SO4, the peak position of which was related to the concentration
of H2SO4 (charge carriers) varying from 0M to 5M. In our case, the only change is the water content
in SPEKK foam. With the loss of water, this relaxation peak shifts to the low frequency range
around 103 Hz, and also gets broader and higher (high loss), suggesting the slower and more
complicated relaxation mechanisms. Unlike the -SO3H- groups bonding with the SPEKK back, the
hydrogen ions are free, therefore, the orientation and arrangement of hydrogen ion-sulfonic group
dipoles are mainly determined by the mobility of hydrogen ions. Obviously, the mobility of
hydrogen ions could be improved with increasing water content, and vice versa. However, the
dissipation factor still is maintained at a rather low level, i.e. it just shows one order increase,
compared with other materials possessing similar permittivity, as shown in Table 1. Lastly, another
important structure feature also accounts for the low dielectric loss, which is the porous structure.
As-prepared SPEKK foam
Figure 8 Comparison of relative permittivity (A) and ε’/tan(δ)
(B) of As- prepared SPEKK foam with literatures
Since the dielectric loss of air is extremely low, the ca.64vol% pores (it can be seen as 64vol% air
filled SPEKK “composite”), can obviously lower the loss of the SPEKK foam.
In this study, the homogenous SPEKK foam with very low density has been successfully
fabricated via an economically efficient and environmental friendly approach, in which, the
concentrated H2SO4 did not only sulfonate PEKK, but also provided sufficient thermal energy
through hydration reactions to generate water vapor as blowing agent. The whole foaming process
is accomplished at low temperature and pressure without handling toxic chemicals, suggesting
unique advantages compared with other conventional foaming techniques with high temperature
and high pressure requirements. The resulting SPEKK foam has low mass density 0.42g/cm3
uniform pore sizes with a diameter of ca. 5μm. Considering the importance of sulfonated aromatic
polymers in energy applications, this approach provides a new and facile way to fabricate
sulfonated aromatic polymers with low density.
The dielectric properties revealed that the SPEKK foam showed great potential as the new
generation of energy storage materials due to the high energy storage ability (dramatically high
permittivity 109 at 0.01 Hz) and energy efficiency (low loss, 2 at 0.01 Hz), In addition, the SPEKK
foam can also maintain the advantages of the polymers without sacrificing the flexibility and low
mass density, unlike polymer composites filled with rigid conductive fillers and ceramics. Moreover,
since the common interests in SPEKK are still limited in its capability as proton exchange
membrane, our findings would definitely extend the applications of SPEKK in energy field.
The authors gratefully acknowledge Cytec Industries Inc. (Cytec Engineered Materials) for
providing PEKK material.
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High Modulus Aliphatic Polyimide from 1, 3-Diaminopropane and
Ethylenediaminetetraacetic Dianhydride: Water Soluble to Self-Patterning
Polymer, 2011, 52, 5186.
In this study, a fully aliphatic polyimide (FAPI) with ultra high modulus and hardness has been
synthesized via 1, 3-diaminopropane and ethylenediaminetetraacetic dianhydride. The thermal
gravimetric analysis (TGA) and differential scanning calorimetry (DSC) studies revealed the
thermal imidization process occurred at a low temperature, around 140~150°C. The resulting
aliphatic polyimide could not only quickly dissolve unconventional organic solvents including
methylene chloride, N, N–dimethylformamide and N-methylpyrrolidone, but also in water at room
temperature, which has not been reported before. Nanoindentation technique applied to evaluate
mechanical properties of the synthesized FAPI indicated that the modulus (Er) and hardness (H) of
FAPI was higher than 8 GPa and 0.3 GPa, respectively, which are approximately 200% of
commercial polyimide products. According to Fourier transform infrared spectroscopy (FTIR)
analysis of FAPI and structure analysis of the two monomers, the existence of nucleophilic
trimethylamine and its interaction with electrophilic carbonyl groups were considered primarily
accountable for both water solubility and high mechanical performances of FAPI. Higher
imidization temperature (250°C) resulted in the self-patterning of FAPI film, which was probably
related to the hydrogen bonding induced assembling of poly(amic acid) as well as the cross-linking
of FAPI at high temperature.
High performance polyimide (PI) resins possess versatile applications in aerospace,
transportation, microelectronics, photonics, etc, due to their outstanding thermal/thermo-oxidative
stability, high mechanical performances, as well as superior electrical and dielectric properties [1-
13]. Thus, the development in past decades of novel polyimides with improved properties and
better processiblity has been of great interest, both commercially and academically. In particular,
high thermal stability [1,3,6,11,12,13] and low dielectric constant [1,2, 5,10, 11] have been the
goals in most studies, while the mechanical performance, such as modulus and hardness, of PIs did
not show significant improvement via synthesis approaches, compared with fabrication of PI
On the other hand, synthesis of highly soluble PIs is also of great significance in practical
applications, such as microelectronics and gas separation which usually require PIs in the form of
thin films. Many studies have revealed that the solubility of PIs could be improved by introducing
bulky and unsymmetrical groups, flexible bonds, large pendent or polar substituents into the PI
chains, using nonplanar/alicyclic monomers [1,2,6]. The resulting PIs showed great solubility in
various organic solvents, including N,N –dimethylformamide (DMF), N- methylpyrrolidone (NMP),
N,N- dimethylacetamide (DMAc), dimethyl sulfoxide (DMSO), Dichloromethane (DCM) and
others . The organo-soluble PIs make solution processing of PIs possible, which significantly
benefits the applications of PIs. However, the great amount of toxic and nondisposable organic
solvents used in solution processing is obviously against the trend of green chemistry, not to
mention the high cost of these organic solvents. The increasing needs for environmental friendly
fabrication techniques/working conditions require the development of PIs readily dissolving in
cheap and less toxic solvents. Undoubtedly, water soluble PIs are with great potentials. In addition,
water soluble PIs do not only solve the existing problems of high cost and toxicity of organic
solvents, but might also possess potential applications in drug delivery systems , ultrafiltration
of dissolved heavy metal ions in water , and widespread applications via metal ion-polymer
interactions , which have been studied in other water soluble polymers.
Aromatic structures are usually believed to be the primary reason for the high thermal stability
and mechanical properties of PIs, besides the interactions between polar imides groups. Thus,
among numerous studies on synthesis of PIs, the aromatic monomers (dianhydride and diamine)
dominated the use of aliphatic monomers. Only a small portion of studies involved using aliphatic
monomers to synthesize semi-aromatic PIs [3,5,8] or aliphatic PIs [10,11], while most of them still
had alicyclic structures on the polymer backbones and similar solubility to conventional fully
In this study, the first water soluble PI with a fully aliphatic chain structure has been
successfully synthesized. More significantly, it shows incredibly high modulus and hardness
determined by nano-indentation, dramatically higher than those of commercial PI products and all
other reported synthetic PIs [11, 12]. Before this study, polyimide with simultaneous water
solubility and high mechanical properties had not been previously reported.
2.1.Materials and Synthesis route
In this work 1,3-diaminopropane (DPA) (liquid, purity >99%, Aldrich) and
ethylenediaminetetraacetic dianhydride (EDTA) (yellow powder, purity>98%, Aldrich) (Scheme 1)
were selected to synthesize fully aliphatic polyimide. The synthesis route was described as follows:
DPA was injected into 120 ml 1-methyl-2-pyrrolidinone (NMP) under continuous and vigorous
magnetic stirring for half an hour to ensure the dissolution of DPA in NMP solvent. Equal molar
amount of EDTA was slowly added to the DPA/NMP solution. Only after several hours, could we
observe a significant amount white precipitate appearing in the NMP, indicating the formation of
poly (amic acid) (PAA), the intermediate polymer for polyimide (PI). The whole reaction process
was carried out in an N2 atmosphere, and a cool water bath was also applied to keep the reaction
temperature low. After 48 hours, a turbid PAA/NMP suspension formed containing both suspended
and dissolved PAA. Before collecting PAA via filtration, the precipitant acetone was added to
precipitate dissolved PAA. The PAA product was rinsed a couple of times to remove the residual
NMP molecules for further imidization and characterization. Two fully aliphatic polyimide (FAPI)
products were obtained via thermal imidization for 3hrs at 150 °C (FAPI-150) and 250°C (FAPI-
250) which are very lower than those in the synthesis of aromatic polyimide. The imidization
temperature was determined via thermal gravity analysis of PAA, and more details will be discussed
in following sections.
2.2.1. Thermal Analysis
The thermal analysis was accomplished via both thermal gravimetric analysis (TGA, TA Q600)
and differential scanning calorimetry (DSC, TA Q20). In the TGA study, the weight loss and
thermal stability of PAA and two PIs were examined at 20°C/min ramp rate below 500°C.
Combined with DSC study, which was carried out at 5°C/min ramp rate from 30°C to 300°C, the
imidization temperature was determined.
2.2.2. Intrinsic viscosity
The intrinsic viscosity of FAPI-150 is 0.20 ± 0.01dl/g at 20°C in deionized water. The
measurement was carried on a Cannon ubbelohde viscometer. The intrinsic viscosity of FAPI-150
given here was the average value of 10 measurements.
2.2.3. Solubility analysis
Generally, PI can dissolve in a variety of organic solvents, such as N,N –dimethylformamide
(DMF) N- methylpyrrolidone (NMP) and Dichloromethane (DCM), which is really important to the
solution processing of PI and its composites. Due to the fully aliphatic structures of FAPIs, it is
speculated their solubility may be different from the aromatic ones, thus, the solubility of PAA and
FAPI in water, DMF, NMP and DCM was studied. Moreover, because the solubility of polymers
strongly depends on the cohesive properties of polymers as well the interactions between polymer
and solvent, the study on solubility could provide useful information of inter- and intra-molecular
interactions of polymer chains.
Fourier transform infrared spectroscopy (FTIR) was used to study the molecular structures of
PAA and FAPIs, as well as the inter- and intra- molecular interactions. The dilute PAA/waster
solution was dropped on to calcium fluoride (CaF2) substrate, and dried at 70°C for an hour to get
PAA thin film sample for FTIR characterization. Corresponding FAPIs thin film samples for were
obtained via drying PAA samples at 150 °C and 250 °C for 3hrs, respectively.
2.2.5. Nanoindentation test
The reduced modulus(Er) and hardness (H) of PAA and FAPIs films were studied by a
nanoindentation approach using a Hysitron TI 900 Triboindentor system. A Berkovich tip (three-
sided pyramidal diamond tip) was selected as indentation probe, and the test was peformed under
open loop mode. Before the test, the indentor was calibrated on a standard polycarbonate sample.
The load-depth curves were recorded and used for the calculation of reduced modulus (Er) and
hardness (H) according to Oliver-Pharr’s method . The film samples with a thickness of ca.
25μm were prepared on silicon wafer using the same method for preparation of FTIR sample. In
order to guarantee the accuracy of nanoindetation test, the indentation was only applied to the
smooth surface of films, which was predetermined by atomic force microscopy in TI 900
2.2.6. Morphological study
The surface morphology of FAPI films which was significantly affected by thermal imidization
process was characterized by both scanning electron microscope (SEM, FEI 200) and AFM
(Hysitron TI 900 Triboindentor System).
3. Results and Discussion
3.1.TGA and DSC results
According to the derivative TGA curve (Fig.1) and DSC result (Fig.2), PAA experienced three
stages with increase of temperature, and they are: stage (1) 100 ~ 200 °C, stage (2) 200 ~275°C and
stage (3) ~320 °C. Below 100°C, neither was there mass loss nor remarkable changes in heat flow,
suggesting no chemical reactions occurred. This is the main reason why we dried PAA and FABIs
at 70°C. In stage 1, multiple peak mass loss appeared in the TGA curves, while in the DSC results,
the most intense exothermal peaks were also observed in this temperature range, especially around
140 ~ 150°C, which is believed to result from the imidization of PAA. It is well known that the
functional groups have different chemical and physical activities on different sites of polymers
chains, thus, the multiple mass loss and exothermal peaks could be observed, corresponding to
stepwise imidization reactions for this aliphatic PAA based on EDTA and DPA. The noisy peaks in
the DSC curves still appeared after repeated testing, and it is probably caused by unstable expansion
and evaporation of water vapor from the sealed aluminum DSC pan, which affects the contact
between the pan and the furnace, however, the strongest exothermal peaks around 1500C are so
evident as to suggest the most intense imidization at this temperature range. According to Scheme 1,
when 1 mole EDTA and 1 mole DPA convert into completely imidized PI via polycondensation and
thermal imidization of PAA, 2 mole water molecules will be generated, correspondingly leading to
-100 0 100 200 300 400 500 6000.0
1500C, ~11.2% loss
0 100 200 300 400 500 600 700 800 900 1000
Figure 1 TGA curves of PAA and FAPIs obtained at 150°C and 250 °C
approximately 10.91% weight loss. From the cumulative TGA curve (Fig. 1B), at a ramping rate of
20°C, when the temperature reaches 150°C, the weight loss is approximately 11.2 wt% which
coincides with the calculated value (10.91wt%), indicating the completion of imidization. Therefore,
150 °C was selected for preparing FAPI-150 sample, to further ensure the complete imidization, the
imidization underwent 3hrs. Compared with the TGA curves of PAA and FAPI-150, we can see,
first, the imidization could be completed in the selected conditions; second, before the thermal
decomposition of polymer chains at ca. 320°C (stage 3) according to the thermal gravity analysis in
Fig.1, another peak appeared at around 250 °C, which existed in both TGA curves of PAA and
FAPIs. Correspondingly, moderate exothermal peaks were observed near this temperature in DSC
curve (Fig.2). Therefore, we also chose 250 °C as the second imidization condition of PAA to find
out the possible reason for this peak. The resultant FAPI-250 sample turned out to be chemically
stable before thermal decomposition. Usually, after immidization reaction, the cross-linking
reaction can happen to PIs with increasing temperature  and /or with addition of linear aliphatic
diamines . Thus, it is reasonable to assume it is the cross-linking of FAPIs account for the peaks
on TGA curves and DSC curve around 250 °C, which will be further proved by the solubility test in
section 3.3. But the cross-linking of FAPI is not the subject of this paper, and thus will not be
discussed in detail.
Compared with other aromatic or semi-aliphatic PAAs and PIs [1,3.6.11,12,13], this new PAA
and FAPI have lower imidization and decomposition temperatures, primarily owing to the absence
of phenyl groups. On one hand, the linear aliphatic PAA has a more densely packed chain structures
that could increase the probability of carbonyl groups attacking amino groups, compared with large
phenyl groups contained PAAs; on the other hand, the phenyl groups are highly thermally stable,
different from alkyl chains. However, despite the reduced stability of both FAPIs at high
temperature, they are still thermally stable below 200 °C, which satisfy most high temperature
applications for PIs. At the same time, the rather low imidization temperature and short imidization
time compared with aromatic PIs, the synthesis of FAPIs in this study is remarkably energy
3.2.Fourier Transform Infrared Spectroscopy (FTIR)
120 180 240
Figure 2 DSC curve of PAA
A rather broad and strong peak around 3400 cm-1
appears in the spectrum of PAA. It is broad
enough to overlap –CH2 absorption peaks at ca. 2900 cm-1
. Also, it has the largest peak strength
regarding to other absorption peaks, suggesting a great amount of –OH and –NH groups on PAA
chains. Since both –OH and –NH could act as both hydrogen donor and hydrogen acceptor in
hydrogen bonding, due to the high electronegativity of O and N, thus, it is believed this structural
characteristics will create many hydrogen bonds in PAA. The peak at ~1640cm-1
was assigned to
the C=O bonds in amides, which has also been observed in similar aliphatic PAAs [19, 20]. After
thermal imidization at 150 °C, owing to the ring closure reaction, –NH groups and –COOH groups
turned into the imide group. Correspondingly, the peak at 3400cm-1
became weaker and narrower,
and only the end groups of PAA chains account for the absorption peak at this frequency. At the
same time, the characteristic peak of –C=O shifted to ca. 1740 cm-1
, owing to the remarkable loss of
hydrogen bonds in FAPI-150. The earlier research has already revealed that hydrogen bonding
impacts on the infrared absorption peaks of carbonyl groups [21, 22]. It is also assumed that the
1000 1500 2000 2500 3000 3500 4000
-C-N- Stretch -C=O
Figure 3 FTIR spectrums of PAA and FAPIs (right) and FAPI-150/water solution (0.02g/ml)
hydrogen bonding affected the absorption peak of C-N bonds. According to Scheme 1, we can find
a great number of C-N bonds in the polymer backbones in this study. Besides the C-N bonds in
amino/imide groups, the C-N bonds in tertiary amines from EDTA monomers also exist. However,
there was no distinct characteristic peak of C-N stretching between 1020 cm-1
to 1250 cm-1
observed in PAA, in contrast to a strong and remarkable peak in FAPI-150 at ca. 1242 cm-1
. Due to
the existence of lone electron pairs in N atoms, the amino groups are good hydrogen acceptors to
form hydrogen bonds. The imidization consumed most hydrogen donors, that is, H atoms in –
COOH and –NH groups, correspondingly, the N atoms with lone electron pair would not
necessarily be involved in hydrogen bonds, thus, it is reasonable to correlate the changes in the
infrared absorptions of –C-N bonds with the vanishing of hydrogen bonding. Furthermore, notice
the spectrum of FAPI-250 did not show any peak shift or changes in peak strength, compared with
FAPI-150, suggesting there were no changes in covalence states of FAPI during cross-linking.
Table 1 Solubility of PAA and FAPIs in various solvents
Water DCM DMF NMP
PAA Soluble Insoluble Soluble Soluble
FAPI-150 Soluble Insoluble Soluble Soluble
FAPI-250 Insoluble Insoluble Insoluble Insoluble
For polymer materials, both inter- and intra- molecular interactions determine their solubility in
solvents, which strongly depend on the molecular structures of the polymer chains. The cohesive
energy density, δ, described by Hansen solubility parameters (HSP) as given in equation (1),
constitutes three parts: dispersion bonds (Van der Waals force) (δD), polar force (δP) and hydrogen
bonds (δH) .
2 2 2 2
D P H (1)
In this study, we tested the solubility of PAA and FAPIs in three different solvents: water,
dichloromethane (CH2Cl2, DCM) and dimethylformamide ( (CH3)2NC(O)H), DMF). Water is a
typical protic solvent which can donate a proton (hydrogen), while the other two are typical aprotic
solvents which cannot dissociate protons. Among them, water has the highest polarity, and DCM
has the lowest polarity. According to Table 1, some interesting results have been found. First of all,
both PAA and FAPI-150 could be quickly dissolved in water at room temperature. On the other
hand, they are insoluble in DCM unlike aromatic thermoplastic PI (Extem resin XH1015, sabic).
Usually, PIs are hydrophobic and polar polymers which could be dissolved in DCM, but not in
water. The reverse solubility for FAPI-150 suggests some unique inter- and intra- molecular
interactions that have not been found in any other PIs.
It is easy to understand the solubility of PAA in water, since a great amount of –COOH and –
NH groups on the backbones could form a lot of hydrogen bonds with highly polar and protic water.
This is also one reason why DCM could not dissolve it. The less polar DCM could not either
provide hydrogen donor or acceptor, thus it is hardly to form hydrogen bonds with PAA. On the
other hand, the linear chain structure of PAA favors the dense packing of PAA chains, which could
strengthen both inter- and intra- molecular hydrogen bonds, as well as other inter- and intra-
molecular interactions, due to the absence of steric effect that is common in aromatic polymers.
This makes it more difficult for the diffusion of solvent molecules into PAA. For aromatic PAAs,
in order to improve their solubility, the tertiary amines were often used to form the water soluble
PAA ammonium salts . In our study, the trimethyleneamino (-(CH2)3N) groups from EDTA
also exist on the polymer backbone. However, owing to the partial solubility of EDTA in water, the
fact of good water solubility of PAAs may suggest that the trimethylamine groups in EDTA are
different from tertiary amines such as triethylamine, and incapable of forming similar PAA
ammonium salt. This further revealed the contribution of hydrogen bonding to the water solubility
of aliphatic PAA.
As discussed above, the imidization results in the disappearance of hydrogen bonds in FAPIs.
However, the FAPI-150 can still dissolve in water, which should be related to the nucleophilic
trimethyleneamino (-(CH2)3N) groups on the backbone. The lone electron pairs on N atoms in
trimethyleneamino group could act as a good hydrogen acceptor which could be easily “attacked”
by water molecules. At the same time, the nucelophilic (electron giving) trimethyleneamino group
could have electrostatic interaction with electrophilic (electron withdrawing) carbonyl groups in
imide. At the same time, the water solubility of FAPI-150 also to some extent came from the
protonation of trimethyleneamino groups , as shown in Scheme 1. Taking into consideration the
trimethylamine groups, such electrostatic interactions could be strong enough to stop DCM. It
should be pointed out here, that although the N atoms in imide groups also have lone electron pairs,
they are delocalized by the resonance of electron-withdrawing carbonyl groups next to it. Therefore,
the conventional aromatic PIs could not form sufficient hydrogen bonds with water and have good
water resistance. DMF also could dissolve both PAA and FAPI-150. On one hand, the structure of
DMF determines it is a good hydrogen acceptor to bond with –COOH and –NH groups; on the other
hand, its high polarity also makes it a good solvent for both PAA and FAPI-150.
For FAPI- 250, there was no solvent capable of dissolving it, which may have resulted from two
possibilities: formation of new and super strong intermolecular interactions, and cross-linking.
According to the FTIR analysis, FAPI-250 has similar covalence states as FAPI-150, which implies
there were no new functional groups forming, such as nitriles ( C N ), thus the inter- and
intra-molecular interactions could not have produced significant changes. Therefore, the resistance
of FAPI-250 to all solvents is believed a consequence of cross-linking at 250°C.
Nano-indentation is an insightful testing technique to investigate the mechanical properties of
matierials. In particular, this technique has been frequently used in microelectronics and thin films.
0 200 400 600 800 1000 1200 1400 16000
Aromatic thermoplastic PI
Figure 4 Load vs. Depth curves of FAPIs and aromatic thermoplastic PI
Polyimide, as an important high performance polymer in these two fields has rarely been studied via
nano-indentation technique. Fig.4 shows the load-depth curves of FAPIs and aromatic thermoplastic
PI obtained by nano-indetation test. Each test includes 5 cycles with increasing indention depth,
and every cycle has three stages: loading, holding and unloading. Obviously, at the same applied
load level, the indention depth on FAPIs are much shallower than aromatic thermoplastic PI (ETEM
resin XH-1005, Sabic), suggesting that FAPIs synthesized in this study are more rigid. Table 2
gives the reduced modulus (Er) and hardness (H) calculated according to Fig.4 and equations (2)
2 21 11 i s
r i sE E E
Where, Ei and Es are the Young’s moduli of indentor material (diamond in this study) and
indented materials (PAA, FAPIs and commercial PI); vi and vs are Poisson’s ratios for indentor and
indented materials. Pmax is the maximum load during indentation, and A is the indentation area.
The commercial product of aromatic thermoplastic PI shows a reduced modulus of 4.36 Gpa
which is only 50% of FAPI-150 (8.92 Gpa) synthesized in this study, so is the hardness: 0.18 Gpa
for aromatic thermoplastic PI vs. 0.31 Gpa for FAPI-150. The highest reduced modulus and
hardness were found in FAPI-250 (Er: 13.50Gpa, H: 0.57Gpa). The extraordinary reduced modulus
and Hardness of FAPI-150 further proved the possibility of the strong electrostatic interactions
between nucleophilic trimethylamine groups and electrophilic carbonyl groups. In aromatic
thermoplastic PI, the high modulus is basically from the polar imide groups and rigid phenyl groups.
In FAPI, owing to the absence of aromatic structures, the inter- and intra-molecular interactions
play more important roles in mechanical properties. Usually, the linear aliphatic polymers have low
rigidity and modulus because of the flexible and compliant chain structures, thus, the very high
modulus should be from the strong inter- and intra-molecular interactions, as mentioned earlier,
which extremely restrict the mobility of polymer chains and improve their resistance to deformation.
This assumption could be further proved by the nano-indentation study of PAA. Both reduced
modulus and hardness of PAA are even higher than that of FAPI-150. This is because, besides the
interaction between trimethylamins groups and carbonyl groups, a great amount of hydrogen bonds
also account for strong inter- and intra-molecular interactions, due to the existence of –COOH and –
NH groups. On the other hand, charge transfer phenomenon has been demonstrated in aromatic PIs
. It was proposed the charge transfer could take place between amine phenyl fragment (electron
donors) and diimide fragment (electron acceptor), which was believed to be the reason for the
amber color of PIs [1, 26]. In this study, due to the similar amber color of FAPIs, the charge transfer
may also happen between carbonyl groups (eletrophilic electron acceptor) and trimethylamine
(nucleophilic electron donor), as an addition to the electrostatic interactions responsible for the high
mechanical performances of PAA and FAPIs. For FAPI-250, besides the reasons above, the cross
linking structures also partially account for the highest Er and H.
An interesting study by humidifying FAPI-150 (FAPI-150-Humidified) provides more proof of
the contribution of inter- and intra-molecular interactions to the great mechanical performance. In
this process, the dry FAPI-150 thin film was exposed to water vapor from a humidifier for about 1
second, resulting in FAPI-150-humidified sample. The load-depth curve of FAPI-150 humidified
sample was given in Fig.5, and its Er and H were listed in Table 2. Its load-depth curve was
significantly changed after being humidified. First of all, at holding stage, the load could not
remain constant as other samples did; secondly, remarkable hysteresis loops formed between two
cycles; thirdly, with very small applied load, the indentation depth is rather huge, compared with
FAPI-150; thirdly, at unloading stage, the minimum load was preprogrammed at non –zero values,
as shown in Fig. 4; however, for FAPI-150-humidified sample, the minimum load went done to
below zero, showing some characteristics of adhesives . The adhesion behaviors of materials
have been studied by indentation approaches . The pull-off force as indicated in Fig.5 was
normally seen as the indicator of adhesion properties. For a good adhesive, the negative pull-off
force should be reached, as a result of the adhesive sticking to the indentor probes. Thus, we can
conclude the FAPI-150-humidified sample was soft, viscous and sticky.
Furthermore, the Er and H of FAPI-150- Humidified sample was extremely low, only 0.17 Gpa
and 9 ×10-4
Gpa, respectively, corresponding to 1.91 % and 0.3% of those of FAPI-150. The
dramatic decline of mechanical performances reflects the breakdown of inter-and intra-molecular
interactions. Because of the same reason introduced in section 3.3 on solubility, the dry FAPI-150
sample will absorb water molecules when exposing to water vapor, resulting in the formation of
hydrogen bonds between water molecules and FAPI-150. The existence of water molecules does
500 1000 1500 2000 2500 3000
0 500 1000 1500 2000 2500 3000 3500
Figure 5 Load vs. depth curve of FAPI-150-Humidified
not only break the inter- and intra-molecular interactions via forming hydrogen bonds with FAPI-
150, also leads to negative steric effect and enlarge the intermolecular distance. Both of them would
cause the changes of the mechanical performances as shown in Fig.5. In other word, the mechanical
properties of FAPI-150- humidified sample further demonstrates the significance of inter-and intra-
molecular interaction among trimethyleneamino, carbonyl, and imide groups, as discussed earlier.
The same degradation of mechanical performances also happened in humidified PAA, but did not
happen to FAPI-250 (no water absorption).
Table 2 Reduced modulus and hardness of PAA, FAPIs and Aromatic thermoplastic PI
PAA 10.51 0.31 0.36 0.02
FAPI-150 8.92 0.45 0.31 0.02
FAPI-250 13.50 0.49 0.57 0.07
4.36 0.05 0.18 0.01
0.17 0.05 0.0009 0.0007
3.5.Self patterning of FAPI-250 sample
The ordering of aromatic PAA chains during drying has been reported . It is well known
that the ordered structures have better mechanical properties, compared with randomly ordered
structures. This structure ordering will be kept when PAA is turned into PI in solid state. It has also
been proposed that the stronger hydrogen bonding prefers near linear geometries . In this study,
it will be easier for the linear aliphatic PAA to assemble into highly ordered linear structures as a
result of plenty of hydrogen bonds, which would exist in both FAPIs. According to the optical
microscopic images in Fig. 6, the film sample of FAPI-150 shows a rather smooth surface, as
indicated by a homogeneous grey color (Fig.6A). However, in FAPI -250, an interesting pattern
(Fig.6B) with local ordering has been observed. The SEM image (Fig.7A) reveals that this pattern
has continuous ridge-like structures, and could be divided into many sub-blocks, in which every
stripe (ridge) is nearly linear and quasi-parallel to each other with a nearly constant inter-distance of
ca. 5um between two ridges. Furthermore, the AFM image (Fig.7B) also revealed a constant height
of those quasi-parallel ridges. Since there was no external field applied in the whole process, this
highly ordered structure, which was not observed in FAPI-150, should result from both ordered
linear PAA chains as a result of hydrogen bonds and high temperature imidziation (cross-linking).
At the same time, because of the fully aliphatic chain structures of FAPIs, there was no
aromatic/aliphatic micro-phase separation which partially accounted for the self-assembled
bisamides . Thus, the high performances of FAPIs in this study also more or less benefit from
the ordered polymer structures. This highly ordered structures came from a self-patterning process
during drying and imidization of PAA, suggesting potentials of FAPIs in micro- and nano-
patterning for microelectronic applications. At the same time, the high modulus and hardness of
FAPI-250 guarantee good dimensional stability of the patterns. The formation mechanisms and
control of the self-patterning will be studied later.
Via polycondensation and thermal imidization at low temperature, 1, 3-diaminopropane and
ethylenediaminetetraacetic dianhydride reacted into the first water soluble aliphatic polyimide,
which is not only benign to the production and environment, but also extends potential applications
of polyimide materials in a water environment, such as bio/med applications, water purification, etc.,
because of the highly hydrophilic trimethylamine groups. At the same time, the strong molecular
interactions between trimethylamine and carbonyl groups, as well as the dense packing of FAPI
chains (weak steric effect) determined the high modulus and hardness, dramatically surpassing all
reported and commercial PI products, i.e. modulus and hardness of FAPI are approximately 200%
of commercial polyimide products. This new FAPI with both improved processibility (water
solubility) and performance/properties will enrich the family of high performance and functional
polyimide materials and extend their usefulness for commercial applications. Furthermore, via
controlling imidization temperature, this aliphatic polyimide can convert into crosslinked structures
Figure 6 Optical images of FAPI-150 (A) and FAPI-250 (B)
showing excellent chemical resistance to all solvents. At the same time, the thermal crosslinking
process can realize well regulated micro-patterns which may benefit the applications in electronic
The authors gratefully acknowledge Sabic Innovative Plastics Co. for providing polyimide. The
authors also gratefully acknowledge Dr. David F. Bahr and Miss Anqi Qiu from School of
Mechanical and Materials Engineering for assistance on nanoindentation tests.
Figure 7 SEM and AFM images of surface morphologies of FAPI-250 film
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