Post on 15-May-2023
Some aspects of preparation methods and properties of polyaniline
blends and composites with organic polymers
Alexander Puda,*, Nikolay Ogurtsova, Alexander Korzhenkob, Galina Shapovala
aInstitute of Bioorganic Chemistry and Petrochemistry, Ukrainian Academy of Sciences, 50 Kharkovskoye Shosse, 02160 Kiev, UkrainebATOFINA, CERDATO, 27470 Serquigny, France
Received 21 April 2003; revised 8 August 2003; accepted 21 August 2003
Abstract
Interest in applications for polyaniline (PANI) has motivated investigators to study its mechanical properties, the
thermostability of its conductivity, its processibility, etc. and its use in polymer composites or blends with common polymers.
As a result, several methods to produce composites/blends containing PANI have been developed, allowing the preparation of a
wide spectrum of such materials. Here, generalized approaches for the preparation of such materials are reviewed. Specifically,
we consider two distinct groups of synthetic methods based on aniline polymerization either (1) in the presence of or inside a
matrix polymer or (2) the blending of a previously prepared PANI with a matrix polymer. Some aspects of these methods are
analyzed, emphasizing features that determine properties of the final composites/blends.
q 2003 Elsevier Ltd. All rights reserved.
Keywords: Polyaniline; Composites; Blends; Preparation method; Properties
Contents
1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1702
2. Synthetic methods to prepare PANI blends and composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1704
2.1. Composites produced by polymerization of aniline in dispersion systems . . . . . . . . . . . . . . . . . .1705
2.2. Chemical in situ polymerization of aniline in the presence of a polymer matrix . . . . . . . . . . . . .1708
2.2.1. Solution polymerization method . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1708
2.2.2. Chemical aniline polymerization in/on solid polymer matrix . . . . . . . . . . . . . . . . . . . . .1710
2.2.3. Electrochemical polymerization of aniline in a matrix . . . . . . . . . . . . . . . . . . . . . . . . . .1712
2.3. Polymer grafting to a PANI surface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1714
3. Blending methods. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1714
3.1. Solution blending. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1714
3.1.1. Blends of substituted PANI. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1715
3.1.2. Blends of soluble aniline copolymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1718
3.1.3. Blends prepared due to counterion-induced solubility of PANI . . . . . . . . . . . . . . . . . . . .1719
3.1.4. Preparation of PANI blends from solutions in concentrated acids . . . . . . . . . . . . . . . . . .1725
3.1.5. Blends prepared of joint PANI base and common polymer solutions in NMP . . . . . . . . .1726
0079-6700/03/$ - see front matter q 2003 Elsevier Ltd. All rights reserved.
doi:10.1016/j.progpolymsci.2003.08.001
Prog. Polym. Sci. 28 (2003) 1701–1753
www.elsevier.com/locate/ppolysci
* Corresponding author. Tel.: þ380-44-559-70-63; fax: þ380-44-573-25-52.
E-mail address: echoc@mail.kar.net (A. Pud).
3.2. Thermally processible PANI blends and composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1727
3.2.1. Composites with infusible PANI . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1728
3.2.2. Polymer blends and composites with fusible PANI . . . . . . . . . . . . . . . . . . . . . . . . . . . .1734
3.2.3. Temperature effects and ageing of doped PANI and its composites. . . . . . . . . . . . . . . . .1739
4. Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1744
References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1744
1. Introduction
The 1977 paper by Shirakawa et al. [1] on
polyacetylene was seminal to the development of
contemporary studies on intrinsically conducting
polymers, ICPs. Since then, the interest in ICPs has
developed through three stages: (1) an initial
interest motivated by their unique properties and
practical possibilities; (2) a decline in interest
owing to difficulties in processing and poor
mechanical properties; (3) renewed interest follow-
ing the discovery of solution and melt processi-
bility of PANI in the early 1990s [2–7]. In recent
years, there has been some optimism that striking
advances in understanding the chemistry and
physics of ICPs [8] will support the development
of large-scale applications, witnessed by the award
of the Nobel Prize in Chemistry to Heeger,
MacDiarmid and Shirakawa in 2000. This trend
has also been recognized by manufacturers, who
are actively investing in research and development
in this field. For example, some leading companies
have discussed their strategy and advances in
applications of ICPs at two European events:
‘Commercializing Conductive Polymers’ in Febru-
ary 2002 and 2003 in Brussels and Barcelona,
respectively.
Although a variety of ICPs have been synthesized
and investigated, polyaniline, polypyrrole, polythio-
phene and their derivatives are most often con-
sidered, due to a good combination of properties,
stability, price, ease of synthesis, treatment, etc. In
some reviews on the subject, one can find analyses
of numerous attempts to apply high conductivity,
electrochromic, catalytic, sensor, redox and other
properties of these polymers to different practical
needs [8–26]. However, since 1984 efforts have
shifted to their use as conducting polymer compo-
sites or blends with common polymers [9,14,
27–31]. This trend has been driven by the need to
replace traditional inorganic conducting fillers and
to improve the processibility of conducting poly-
mers, along with their mechanical properties and
stability. These composite materials have introduced
conducting polymers to practical applications in
different fields, including electromagnetic shielding
and microwave absorption [25,32–34], static elec-
tricity dissipation [35–37], heating elements (cloth-
ing, wall papers, etc.) [38,39], conducting glues
[40], conducting membrane materials [41,42], paint
coatings for anticorrosion protection [43], and
sensor materials [44,45].
Among ICPs, PANI is known as having
probably the best combination of stability, conduc-
tivity and low cost [2,18,46]. As a consequence, its
conducting composites are very close to appli-
cations on a large scale for the industrial
applications mentioned above [25,32–45]. Never-
theless, the choice of the best method to produce
composites with specified characteristics remains an
unresolved problem. The problem arises because
the processing method may significantly determine
the properties of the manufactured composite
materials. Known methods to produce PANI
containing composites [31] may be essentially
reduced to two distinct groups: (1) synthetic
methods based on aniline polymerization in the
presence of or inside a matrix polymer, and (2)
blending methods to mix a previously prepared
PANI with a matrix polymer. Roughly, these
include:
(1) Synthetic methods
† Dispersion polymerization of aniline in the pre-
sence of a matrix polymer in a disperse or
continuous phase of a dispersion;
† Chemical in situ polymerization of aniline in a
matrix or in a solution with a matrix polymer;
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531702
† Electrochemical polymerization of aniline in a
matrix covering an anode;
† Polymer grafting to a PANI surface;
† Copolymerization of aniline with other monomers
resulting in the formation of soluble aniline
copolymers, which can be considered as a
composite polymer.
(2) Blending methods
† Solution blending soluble matrix polymers and
substituted polyanilines;
† Solution blending soluble matrix polymers and
PANI doped by functionalized protonic acids
(counterion-induced processibility);
† Solution blending undoped PANI with polymers
soluble in amide or acidic solvents
† Dry blending followed by melt processing (MP)
(mechanical mixing of doped PANI with thermo-
plastic polymer, then molded in a hot press or
extruder);
Naturally, each of these methods has its own
advantages and limitations. Specifically, the synthetic
Nomenclature
ABS acrylonitrile–butadiene–styrene copoly-
mer
AMPSA 2-acrylamido-2-methyl-1-propanesulpho-
nic acid
APS ammonium persulfate
BEHP bis(2-ethylhexyl)hydrogenphosphate
CoPA copolyamide 6/6.9, a random copolymer of
51% [HN–(CH2)5–CO] and 49% [HN–
(CH2)6–NH–CO–(CH2)7–CO]
CSA camphorsulfonic acid
DBSA dodecylbenzenesulfonic acid
DEHEPSA di(2-ethylhexyl)ester of phthalosulfonic
acid
DiOHP di-i-octyl phosphate
DMF N,N0-dimethylformamide
DOP dioctyl phthalate
DPHP diphenyl phosphate
EB emeraldine base
EPDM poly(ethylene-co-propylene-co-diene-
monomer)
ICPs intrinsically conducting polymers
LDPE low-density polyethylene
LEB leucoemeraldine base
LG lauryl gallate
LLDPE linear low-density polyethylene
MSA methanesulfonic acid
NMP N-methyl-2-pyrrolidinone
PA polyamide
PAM polyacrylamide
PANI polyaniline
PC polycarbonate
PCL poly-1-caprolactonePEO poly(ethylene oxide)
PET poly(ethylene terephthalate)
PETG poly(ethylene terethphalateglycol)
PMMA poly(methyl methacrylate)
PMT poly(m-toluidine)
POMA poly(o-methoxyaniline)
POT poly(o-toluidine)
PPD-T poly( para-phenylenediamine)terephthalic
acid
PS polystyrene
PSS poly(styrenesulfonate)
PU polyurethane
PVA poly(vinyl alcohol)
PVDF poly(vinylidene fluoride)
PVC poly(vinyl chloride)
SBS styrene–butadiene–styrene
SPDA sulfonic acid of 3-pentadecylanisole
SPDP sulfonic acid of 3-pentadecylphenol
SPDPAA sulfonic acid of 3-pentadecylphenoxy
acetic acid
THF tetrahydrofuran
TSA p-toluenesulfonic acid
UHMW-PE ultra high molecular weight polyethy-
lene
DSC differential scanning calorimetry
DTA differential thermal analysis
EMI electromagnetic interference
ESR electron spin resonance
TGA thermogravimetric analysis
XPS X-ray photoelectron spectroscopy
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1703
direction is probably preferable if it is necessary to
produce inexpensive conducting composites, due to
use of inexpensive aniline instead of more expensive
PANI, or when there is a need to form composites
which have conductivity only in a thin surface layer.
Good homogeneity and a low percolation threshold
characterize these composites. On the other hand,
blending methods sometimes seem to be more
technological desirable from the standpoint of large-
scale production, particularly in the case of melt
procession techniques. Blending methods will be
probably become very practicable when techniques to
produce inexpensive, nanosized PANI are well
developed.
This review will survey both of the above-
mentioned methods, and the results of studies of the
resulting PANI composites to elucidate the appli-
cation of each method. However, it should be noted
that it was impossible to include here all of the
publications on the topic because of their vast
number, which is, in fact, increasing with every
issue of the specialized scientific journals. As a
consequence we tried to consider publications,
which from our point of view, illustrate the main
aspects of PANI composites.
2. Synthetic methods to prepare PANI blends
and composites
Aniline polymerization in acidic medium results
in the formation of a protonated, partially oxidized
form of PANI [16,47]. This process is sufficiently
complicated to be considered as a specific kind of
cationic polymerization [16]. During the polymeriz-
ation, the PANI chain propagation terminates with
the formation of the most conductive PANI form,
the emeraldine oxidation state, which may be
converted to the corresponding EB by treatment
with an alkali solution, or by rinsing with an excess
of water [16,48].
It was discovered that non-conductive PANI
may exist in a continuum of oxidation states,
changing from the completely reduced leucoemer-
aldine ðy ¼ 1Þ; through the EB ðy ¼ 0:5Þ; up
to the completely oxidized pernigraniline ðy ¼ 0Þ
[16,48,49]:
ð1Þ
Imine sites of the intermediate PANI base
forms are easily protonated, with a striking
insulator–conductor transition, induced due to the
appearance of positive charges in the lattice, while
the number of p-electrons remains constant. As a
consequence, new optical, conductive and para-
magnetic properties appear in doped PANI,
specifically in emeraldine ðy ¼ 0:5Þ salt, for
which polaronic lattice structure I was proposed
by Stafstrom et al. [50]:
ð2ÞObviously, this is an ideal structure, perhaps
realized under ideal conditions (e.g. PANI in its
emeraldine form, protonation effected in dilute
PANI solution). In the solid phase the protonation
of PANI or its composites is limited by diffusion
of the dopant (acid) to imine sites, a process that
may depend on the dopant anion size and the
polymer matrix morphology. As a result, a
homogeneous redistribution of polarons along
PANI macromolecular chains is possible in small
clusters, differing in size, as determined by the
packing of the macromolecules in the material.
Probably, this can be easily checked by measuring
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531704
the conductivity of the same emeraldine form or
its composite sample redoped in the solid phase
condition by acids with large anions.
On the other hand PANI conductivity properties
are a function of not only of the degree of protonation
and oxidation, but also of structural and confor-
mational factors, which may be affected by aniline
polymerization conditions [16,47–50]. This means
that one of the important tasks in PANI synthetic
chemistry is the development of technological
methods leading to conducting composites containing
doped PANI with the best combination of these
parameters. In turn, these properties should match
with other requirements for the final composite
materials.
2.1. Composites produced by polymerization
of aniline in dispersion systems
This method, based on experience in aniline
polymerization, is conducted at low temperatures
(typically 0–5 8C) using an appropriate oxidant
(usually APS, but sometimes K2S2O8, KJO3, H2O2,
etc.) in the presence of water soluble polymers or
tailor-made reactive copolymers [51] (e.g. poly(2-
vinylpyridine-co-p-aminostyrene) [52], PVA [53,54],
poly(N-vinylpyrrolidone) [55,56], PEO [57], cellulose
derivatives [58,59], poly(methylvinylether) [60],
etc.). The technique results in sterically stabilized
colloidal dispersions of PANI particles of different
size (typically from tens to hundreds of nanometers)
and morphology. These colloids can be further mixed
with film-forming latex particles or with stable matrix
polymer dispersions to produce conducting compo-
sites [55,57,61]. Thus, Banerjee and Mandal [62]
synthesized a dispersion of non-spherical PANI
particles with diameters of 150–300 nm, stabilized
with poly(methylvinylether). These particles were
disintegrated into nanosized particles with diameters
less than 20 nm, which were used to prepare
conducting blends with conventional polymers PVC,
PS, PMMA, poly(vinylacetate) and PVA by sonicat-
ing a suspension of the preformed submicronic PANI-
HCl particles in solutions of the matrix polymers. The
blend films exhibited an extremely low percolation
threshold ðfpÞ in every case, with a volume fraction of
PANI-HCl at the percolation threshold in the range of
2.5 £ 1024–4 £ 1024 vol%. The PANI-HCl/PVA
films exhibited self-assembly of nanoparticles of
PANI-HCl. The network was fibrillar, in contrast to
the globular network found with PANI-HCl/PVC
[63]. This difference in morphologies might
arise from differing thermodynamic interactions
between PANI-HCl and the matrix polymers. PVA
was reported to have some affinity with PANI through
hydrogen bond interaction. This affinity may result
in finer dispersions in PVA, and a fibrillar
morphology.
A similar technique [62,63] was used by Beadle
et al. [64] in the polymerization of aniline in the
presence of a film-forming chlorinated copolymer
latex.
Comparatively, small-molecule surfactants were
used for stabilization of ICPs colloids [51]. This
was demonstrated in 1993 by DeArmitt and Armes
[65] in a polypyrrole dispersion produced in the
presence of sodium dodecylbenzenesulfonate. In
this case the surface of the polypyrrole particles
was enriched with the surfactant [51]. The
polymerization of aniline inside micelles of sodium
dodecylsulfate produced a reasonably stable colloid
containing low molecular PANI. This PANI anionic
micellar system had a metal–insulator transition
from the emeraldine salt to EB at the unexpectedly
high pH of 7–8 [66].
Ruckenstein et al. [67–70] have developed
emulsion pathways for the preparation of conductive
PANI composites using the stabilization of an
emulsion by a surfactant. Specifically, they reported
a method to produce PANI/PMMA [67] and PANI/
PS [68] composites via an oxidative aniline polym-
erization carried out by adding an aqueous solution
of the oxidant (APS) and dopant (hydrochloric acid)
to a concentrated emulsion containing an aqueous
solution of the ionic surfactant (sodium dodecylsul-
fate) as the continuous phase and an organic
(benzene) solution of the host polymer and aniline
as the dispersed phase. The corresponding compo-
sites were obtained by co-precipitation of the host
polymer and PANI, with a percolation threshold of
,2–10 vol% PANI. Later, Ruckenstein et al. [69]
developed an inverted emulsion pathways to prepare
PANI composites with SBS rubber at different molar
ratios of aniline/dopant (sulfonic acids), oxidant/
aniline, quantities of a surfactant and nature of the
solvent in the continuous phase. These changes in
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1705
the reaction mixture affected the conductivity and
mechanical properties of the final composite (pressed
at 150 8C), and produced a composite with a
percolation threshold of ,6.9 wt% PANI, and a
tensile strength 7 MPa. In the case of TSA as the
dopant, the composite with 24.6 wt% PANI had a
conductivity as high as 2.5 S/cm.
As far back as 1987 Yassar et al. [71] reported an
alternative method to produce conducting colloid
latexes, through pyrrole emulsion polymerization in
sulfonated and carboxylated PS latexes, in which the
particles were overcoated by polypyrrole. Wiersma
et al. [72] have shown that a critical condition for
stability of such latexes (e.g. PU) is the presence at
latex particles of a chemically grafted non-ionic
polymeric stabilizer, such as PEO or hydroxymethyl-
cellulose. They used transmission electron
microscopy to reveal a ‘core–shell’ morphology of
latex particles (core) coated by the conducting
polymer (shell). These coated particles displayed the
good film-forming properties of the parent PU at
ambient temperature, despite the fact that the low Tglatex component was encapsulated within the high Tgconducting polymer. The composite films produced
had a conductivity in the range of 1025–101 S/cm
[72]. It should be noted that unlike the relatively
smooth and uniform morphology of polypyrrole
coated latex particles [73,74], PANI overlayers
(core) on latex (PS) particles (shell) are rather
inhomogeneous [75–77]. Armes et al. [75] used
XPS to examine the surface compositions of the PANI
overlayers deposited onto micrometer-sized poly(N-
vinylpyrrolidone)-stabilized PS latex particles under
various synthesis conditions for seven preparations.
The thickness of the PANI overlayer was in the range
of 2–30 nm, and the conductivity of the coated
particles substantially increased with a raise of PANI
loading to attain a maximum conductivity 0.17 S/cm
for 9.3 wt% PANI. It has been shown that relatively
rapid polymerization at room temperature resulted in
the non-uniform PANI coatings and reduced PANI
surface yield: Non-uniform PANI coatings were
obtained for the polymerization of aniline hydrochlo-
ride in the presence of HCl in the latex medium at
ambient temperature (25 8C), but more homogeneous
PANI coatings were obtained at 0 8C. The maximum
PANI coverage was found to be around 57–59%,
which is much lower than the surface composition of
94–100% found for polypyrrole deposited onto a
similar micrometer-sized PS latex [78]. Finally, the
improved uniformity of the PANI overlayers prepared
using aniline hydrochoride in the absence of HCl is
consistent with the higher coalescence temperature
found for these PANI-coated PS particles in hot-stage
optical microscopy studies.
The formation of electrostatically bound anilinium
cations in the emulsion polymerization of aniline in
latexes containing polymer particles with surface
acidic (sulfonic) groups may be the origin of increased
homogeneity of the PANI overlayers observed in
these materials. Kim et al. [79] confirmed this
supposition for the aniline hydrochloride polymeriz-
ation in a PS–PSS latex, reporting that a high
concentration of aniline was needed to coat all the
core particles uniformly because of a very small size
of the PS–PSS core particles (of 30–50 nm in
diameter). The conductivity of the produced compo-
site measured on cold pressed pellets and increased
from 2.6 £ 1025 S/cm at 3.41 wt% PANI to a
maximum of 0.05 S/cm at 12.3 wt% PANI. In some
cases it is important to produce a final conducting
composite with good thermostable properties, specifi-
cally for melt processing (MP) techniques. This
suggests that it is preferable to carry out the emulsion
aniline polymerization in latex media in the presence
of sulfonic acids than hydrochloric acid [80] to ensure
higher thermostability of the composite conductivity.
Moreover, the sulfonic acids act both as a surfactant
and as a dopant for PANI [81]. Using this approach,
Xie et al. [82,83] prepared PANI/SBS [82,83] and
PANI/chlorosulfonated polyethylene (CSPE) [84]
composites by aniline polymerization in an emulsion
comprising water and xylene containing the elasto-
mers and DBSA. The composites obtained were
processed by MP or solution processing (SP).
Percolation thresholds were lower for PANI/SBS
(10 wt% for MP sample and 7 wt% for SP sample)
than for PANI/CSPE (14 wt% for MP samples and
22 wt% for SP samples). At the same content of
PANI, the conductivity of the SP composite was
higher than that of the MP composite for PANI/SBS,
with the reverse observed for PANI/CSPE. The
elastomer nature also affected relationships between
mechanical properties and the PANI content, as well
as the morphological structure of the composites.
Thus, for MP samples of PANI/SBS, the composites
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531706
behaved like a thermoplastic elastomer when the
PANI content was lower than 12 wt%, with a high
elongation (about 600%) and low permanent set
(,50%). In the case of PANI/CSPE, a thermoplastic
behavior was observed at higher PANI content
namely between 12 and 18 wt%, with an ultimate
elongation .400% and permanent set ,30%. On
secondary doping of the SP samples with m-cresol,
the conductivity of PANI/SBS increased by two
orders of magnitude and that for PANI/CSPE
increased by six orders of magnitude [82,84]. From
our point of view, these effects indicate that the
interaction of PANI with the elastomers is enhanced
for the more polar CSPE.
The strong effect of interactions of PANI with its
host polymer on the composite properties was
confirmed also by Jeon et al. [85]. They found this
effect for composites of PANI-DBSA/PC prepared by
an inverted emulsion polymerization method devel-
oped in accord with Ruckenstein et al. [67–70]
pathways, in which the role of surfactant and dopant
was played by DBSA [85]. Investigating the effect of
DBSA concentration in the emulsion reaction mixture
on the final composite conductivity, they found that the
electrical conductivity of the composite increased by
about three-fold from a value of 4.5 £ 1023 S/cm (at
16.7% PANI) as the mole ratio of DBSA/aniline was
increased from 0.75 to 3. FTIR spectroscopy on the
composite showed the existence of hydrogen bonding
between PANI and PC, which increased the glass
transition temperature with increasing PANI content.
Moreover, comparison of DSC and conductivity data
showed that the electrical conductivity increased
around the glass transition temperature. The authors
explained this by the fact that the PANI chains
contacted more frequently and facilitated electron
transfer through the hydrogen bonding between PANI
and PC. In addition, the tensile strength of the
composite decreased with PANI content below the
percolation threshold (13 wt%) of PANI (Fig. 1a). This
suggested that PANI functioned as a defect in the PC
matrix in accord with scanning electron microscopy
(SEM) data, which showed an inhomogeneous distri-
bution of PANI in the PC matrix below 13 wt% of
PANI [85]. In contrast, the continuous increase of the
tensile moduli of the composites (Fig. 1b) is attributed
to the higher rigidity of PANI molecules [85]. It is
difficult to demonstrate discrete PANI and PC
components by SEM above the percolation threshold
(13 wt%), suggesting a fine distribution of PANI in the
matrix. Together with the mechanical behavior [85],
this suggests that the structure of the PANI/PC
composite is changed at high content of PANI due to
a physical –chemical interaction (e.g. hydrogen
bonding) of the components. This interaction may
Fig. 1. (a) Tensile strength and (b) tensile modulus of the PC and the
PANI–PC composites as a function of PANI content [85].
Reproduced from Jeon, Kim, Choi and Chung by permission
of Synth Met 1999;104(2):95. q 1999 Elsevier Science Ltd,
Oxford, UK.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1707
also be displayed by improved thermal stability of
PANI/PC blends [86]. Jeevananda et al. [86] used
sodium laurylsulfate (SLS) and TSA, which acted as
the surfactant and as the protonating agent for the
resulting polymer, to prepare these blends and PANI by
one-step emulsion polymerization technique. The
conductivity of the final PANI/PC blends
decreased from 4.70 £ 1022 S/cm (PANI/PC1) to
5.68 £ 1025 S/cm (PANI/PC3) with the change from
TSA to SLS, respectively.
The preceding discussion reveals the importance
of PANI–matrix polymer interactions for the
properties of composites. Such interactions develop
at the aniline polymerization stage, both among the
dispersion (emulsion) particles and at their surface,
and may be also be affected by adsorption of
aniline, acid and oxidant at the surface of the core
particle. As a consequence, their concentration and
physical–chemical interaction with the core particle
surface are important. This may have also a special
significance if the surface contains groups that
interact with these reagents and facilitate their
adsorption to form an adsorbed layer where the
formation of PANI can proceed.
The great number of factors affecting the aniline
polymerization in matrix polymer dispersions and
their impact on the composite properties demands a
strict control of every stage of the polymerization
method. On the other hand it is possible to avoid at
least a part of such complications when using a simple
mixture of a previously prepared nanosized PANI
with a matrix polymer dispersion. Haba et al. [87]
successfully used this approach to produce PANI
containing blends by mixing dilute aqueous disper-
sions (,0.8 wt%) of a nanosized PANI-DBSA with
an aqueous emulsion of the matrix polymer (PMMA
or PS, or a commercial acrylic latex), followed by
water evaporation. The separated powder or mixed
films were then sintered (at 80–120 8C under
pressure), followed by compression molding (at
120–180 8C) of the free samples and fast cooling.
The final blends exhibited an electrical conductivity
of 1026 S/cm at a very low PANI-DBSA content
(0.5 wt%), and tended to plateau above 2 wt%
PANI-DBSA, without a sharp percolation transition.
These results were explained by a significant and fast
segregation process, beginning with the formation of
the PANI-DBSA/polymer aqueous dispersions.
This strong segregation stemmed from the different
surface characteristics of the PANI-DBSA and matrix
polymer particles. The authors emphasized that the
segregation in these systems took place in a very low
viscosity aqueous medium, and was thus very likely a
fast process, in contrast to a segregation phenomenon
in solution cast films, or within a polymer melt. They
found that the conductivity level of the various blends
depended on the PANI content, on the surfactant
present in the polymer matrix emulsion, and was
practically independent of the polymer matrix nature.
The last was accepted as a further proof that the
particle surface characteristics (each polymer particle
was coated with its surfactant) are a key factor in the
segregation process, rather than the character of the
polymer particle itself [87].
2.2. Chemical in situ polymerization of aniline in
the presence of a polymer matrix
Unlike the dispersion systems considered above,
there are other methods of chemical polymerization of
aniline in the presence of polymer matrix which do
not demand the presence of surfactants in the reaction
mixture. Specifically, these are the chemical polym-
erization of aniline by a variety of methods: in a
solution of aniline and a matrix polymer [88,89]; at
the surface of a polymer substrate dipped in aniline
and oxidant solution [90]; directly in a polymeric
matrix, swelled in aniline and contacting with an
oxidant solution [91,92]; or in a polymeric matrix
containing an oxidant and contacting with a solution
or vapors of a monomer [93,94]; etc.
2.2.1. Solution polymerization method
Obviously, it is difficult to find a well-defined
boundary between solution polymerization systems
and the nanosized dispersion methods reviewed
above. This is rather a problem of definitions of
true and colloid polymer solutions, and is a topic for
discussions of PANI containing systems [95,96]. We
may consider that aniline polymerizations follow
from case one to another, dependent on the
polymerization degree and other components of
the system. Specifically, this is characteristic for
water systems containing water-soluble polymers,
e.g. PVA, poly(acrylamide), Nafion, and polysac-
charides [31,89,97,98]. In these systems aniline
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531708
polymerization was usually carried out at lower
temperatures (,0–10 8C), but there were some
differences in the procedure, especially in the
sequence of the addition of reagents to the reaction
mixture. Thus, in some cases aniline was added to
an acidified solution of a matrix polymer (PVA,
chlorinated copolymer latex Haloflex) and oxidant
APS, followed by precipitation and filtration of a
conducting composite [64,88]. In another sequence,
an acidified solution of the oxidant was added to a
previously cooled solution of aniline and polymer,
to effect the polymerization for 3–24 h at lower
temperatures. In particular, Gangopadhyay et al.
[89] used the last approach to prepare a PANI/PVA
composite in an aqueous solution of PVA (1 g in
10 ml water), maintaining the molar ratio of aniline:
APS ¼ 1:1 (at different quantities of aniline), and
pH ¼ 1 in the presence of HCl at ,10 8C. The
polymerization was allowed to proceed for 3 h, and
stopped with the formation of a solution of the
bright green stable composite that could be stored or
precipitated by methanol. In this system, the aniline
polymerization yield was 82–84.5% at low aniline
concentration (0.1–0.2 M). This yield decreased by
up to 52.5% on increase of the aniline concentration
to 0.5 M. These data match those of Stejskal et al.
[56] for an aniline disperse polymerization in the
presence of PVA, for which the PANI yield did not
exceed 40%, albeit at lower concentrations of the
oxidant (APS). The aniline concentrations was
varied from 0.1 to 0.5 M to increase the fraction
of PANI in the composite from 7.75 to 21.01%,
despite the decrease in yield [89]. The final
composite showed good film forming ability with
a conductivity 6.1 £ 1026 S/cm at 7.75% of PANI,
and 1.32 S/cm at 21.01% of PANI. This kind of
conducting PANI composites exhibit significant
EMI shielding capacity, and potential for sensing
moisture and methanol vapor [89,99]. Mechanical
studies show that at moderate PANI content (7.75%)
the tenacity of PANI/PVA composite films decreased
from that of the pure PVA network, probably due to
some disruption of the PVA network, with some
regain on increased PANI loading. The changes at
higher PANI loadings were explained as a direct
consequence of a semi-cross-linked structure of the
matrix polymer, or of a semi-interpenetrating net-
work formed during aniline polymerization [89,100].
But it seems we may accept here an additional
explanation of these changes via a physical–
chemical interaction of PANI and PVA, mentioned
above for the PANI/PC composite, characterized by
similar behavior [85].
A strong (chemical) interaction between PANI
and a soluble matrix polymer can sometimes be
formed due to aniline grafting to radicals appearing
in the polymer matrix backbone under the action of
an oxidant, which in parallel initiates aniline
polymerization in the solution. Obviously, this
possibility depends on the matrix polymer. Xiang
and Xie [101] showed that aniline could be graft
copolymerized onto the backbone of PAM in
aqueous HCl solution in the presence of APS as
oxidant. They dissolved the copolymer PAM-g-
PANI in 5 wt% NaOH solution when the molar ratio
of aniline/acrylamide (An/AM) in the feed compo-
sition was lower than 15. After removal of the salt
ions by dialysis and evaporation of the solution, a
thin film of PAM-g-PAn was obtained and doped by
HCl gas. When the molar ratio of An/AM in the
feed composition was about 15, the HCl doped thin
film of PAM-g-PAn possessed a high conductivity
of 8.8 S/cm.
Ghosh et al. [102] investigated a similar system
and found that PANI synthesized in an optically clear
aqueous solution or dispersion with the support of
aqueous PAM (2–5%) showed excellent storage
stability, due both to limited grafting of PANI on
PAM and to a template effect through hydrogen
bonding between segments of the two polymers.
When investigating the effect of aniline concentration
in the reaction mixture they observed an upper
limiting conversion of nearly 75% [102]. Scanning
electron micrographs showed that a PANI/PAM
composite at low PANI loading (2%) had a little
phase separation, but that a minor phase separation
appeared for a somewhat higher PANI content (10%),
without a gross phase aggregation. The phase
morphology of PANI/PAM composites having even
40% PANI content showed a very intimate and
uniform distribution of the two phases, without the
significant phase aggregation. This highly uniform
phase morphology of the PANI/PAM composites is a
direct consequence of a mutual interaction between
PANI being formed and PAM in the solution during
polymerization of aniline, including the establishment
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1709
of PANI/PAM hydrogen bonding and grafting of
PANI on PAM, as already mentioned. The PANI/
PAM composites showed enhanced thermally stable
electrical conductivities (1028–1021 S/cm) in com-
parison with PANI itself [102].
Water insoluble polymers may also be used to
produce conducting PANI composites through the
solution polymerization method. Specifically, it was
demonstrated by the chemical aniline polymerization
in PS solution in xylene [98]. It was realized by the
addition of the oxidant and DBSA dissolved in xylene
to the xylene solution comprising aniline and PS. The
electrical conductivity of the separated PANI/PS
composites improved with increasing amount of
PANI, to reach a value of 0.1 S/cm at 12 wt%
PANI. Due to the fact that DBSA served as dopant,
these composites were soluble in a variety of organic
solvents (chloroform, xylene, and NMP).
2.2.2. Chemical aniline polymerization in/on solid
polymer matrix
Unlike aniline polymerization in a solution these
methods produce modified polymer matrixes with a
PANI layer at their surface or inside a thin subsurface
layer. Naturally, the thickness and conductivity of the
layers depend on the method of modification and on
the time of contact of the solid matrix with the
reaction medium. These methods produce composites
with a wide surface conductivity range, from semi-
conductor up to the conductivity of pure PANI. Even
a simple dipping method resulted in conductivity of
1–5 S/cm and transmittance of 80% at 450–650 nm
for a 0.5 mm PANI layer [9]. Apparently, this method
is not technological suitable for sheet materials, both
because it requires the use polymer matrixes with a
good adhesion to PANI, and because it produces pure
PANI at the matrix surface, having poor mechanical
properties. At the same time, for fiber and textile
materials with a well developed reactive surface, it
may lead to the production of conducting fibers and
fabrics with grafted PANI at the surface and inside of
pores. This approach resulted in suitable materials for
EMI shielding, sensors, static electricity dissipation,
etc. [9,103,104].
Two methods to obtain electrically conductive
fabrics by in situ polymerization of aniline were
compared by Oh et al. [105]. These materials were
prepared by immersing the Nylon 6 fabrics in pure
aniline or an aqueous hydrochloride solution of
aniline followed, by initiating the successive direct
polymerization in a separate bath (DPSB) or in a
mixed bath (DPMB) of oxidant and dopant solution
with aniline. The authors showed the DPMB process
produced higher conductivity in the composite
fabrics, reaching 0.6 £ 1021 S/cm. Moreover, this
process induced a smaller decrease in the degree of
crystallinity than the DPSB process [105]. In our
opinion, this difference can be connected with the fact
that in the case of DPSB process the Nylon 6 fabrics
was swollen with aniline, which when localizing in
amorphous regions of the matrix acts as plasticizer
and may effect the orientation and arrangement of
Nylon 6 macromolecular segments located there. The
PANI/Nylon 6 composite fabrics displayed a good
serviceability [105]. Thus, no important changes in
the conductivity were observed after abrasion of the
composite fabrics over 50 cycles and multiple acid
and alkali treatment. The stability of the conductivity
decreased by less than one order after exposure to
light for 100 h, but it was significantly decreased after
washing with a detergent [105]. The serviceability of
these materials was improved by plasma treatment of
the Nylon 6 fabrics, resulting in improved adhesion
properties, change of rate of aniline polymerization
[106], conductivity and durability [107].
As in the case of the solution polymerization
method (see above and Refs. [101,102]) the use of
peroxosalts as oxidants causes a graft copolymeriza-
tion of aniline and its derivatives onto a polymer
matrix [108]. Anbarasan et al. [103,104,109] investi-
gated the kinetics of this grafting onto PET, Nylon 66,
wool and Rayon fibers, and proposed a possible
mechanism of graft and homopolymerization of
aniline. Specifically, they carried out oxidative
chemical polymerization of aniline using peroxydi-
sulphate and peroxomonosulphate as the sole initiator
in an aqueous acidic medium in the presence of the
fibers. This resulted in the chemical grafting of PANI
onto the fibers, confirmed by FTIR spectroscopy,
cyclic voltammetry, weight loss study, and conduc-
tivity measurements. The authors proposed a probable
mechanism to explain the experimental results,
describing the graft polymerization of aniline through
interaction of the oxidant with the fiber surface,
inducing the formation of radical sites at the
fiber surface, followed by grafting aniline with its
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531710
subsequent participation in a typical aniline oxidative
polymerization.
Polymerization of aniline on porous materials has
also been used to prepare conducting membrane
materials whose permeability and other properties
could be maintained by the porosity of the final
material and conducting polymer layers formed inside
pores [110,111]. Specifically, Tishchenko et al. [110]
elaborated composite systems based on a microporous
polyethylene membrane modified in situ during the
oxidative polymerization of pyrrole from the gas
phase or by the polymerization of aniline in an
aqueous medium. The composite membranes dis-
played a low resistance in electrolyte solutions owing
to the coating of polypyrrole or PANI inside the pores.
Moreover, according to Elyashevich et al. [111],
analogous composite systems were found with better
thermostability than the parent microporous poly-
ethylene, and demonstrated considerably lower
shrinkage upon heating, probably due to the stiffness
of the conducting polymer coating. A stabilizing
effect of conducting polymers was found even in the
melted composites, in which the oriented state was
maintained on heating samples to temperatures
exceeding the polyethylene melting point by several
tens of degrees [111]. Furthermore, taking into
account the conductivity of these membrane
materials, it may be possible to apply an electric
potential sufficient to control their permeability and
selectivity in solution.
The changeover from polymerization of aniline in
interstices or at the surface of a fiber/textile or a
porous unswollen materials, to its polymerization as
imbibed in a polymeric matrix (using so-called
diffusion-oxidation method [9]), results in the for-
mation of a thin surface/subsurface conducting
composite layer of high transparency [112]. Properties
of such composite materials depend on a physical–
chemical interaction (e.g. hydrogen bonding) between
PANI and the host polymer [92,113,114], and should
also be influenced by interaction of the latter with
aniline formed during swelling [115]. Specifically,
Byun and Im [114] prepared a PANI/Nylon 6
composite by immersing a Nylon 6 film swelled
with aniline in APS solutions containing different
acids (hydrochloric, benzenesulfonic, sulfosalicylic
and TSA). The composites consisted of three layers:
two outer ones were conducting composite layers and
the inner one was pristine Nylon 6. These composites
displayed very low percolation threshold content
(about 4 wt%) and provided a conductivity of about
3.5 £ 1022 S/cm at 4.4 wt% PANI-HCl content.
Hydrogen bonding between PANI and Nylon 6 was
found to affect the doping characteristics of the
composite, to result in a much lower doping level for
the composite than that for pure doped PANI [114].
The strong effect of physical–chemical interactions
among the composite components on its properties
was additionally confirmed by the results of dynamic
mechanical thermal analysis. Specifically, the crystal-
line regions of Nylon 6 in the composites are partly
destroyed by the formation of PANI, as deduced from
the reduced heat of fusion, degree of crystallinity and
melting point of the composite in comparison with the
parent Nylon 6 (Table 1) [114].
The physical–chemical interaction of aniline with
the host matrix polymer not only affects the composite
properties, but also results in high specificity of the
aniline polymerization process inside thematrix. Thus,
Pud et al. [115] found that this process could be run in a
PET matrix only in a chlorine (bromine) containing
water medium under the action of products of the
halogen hydrolysis. Moreover, the polymerization
outcome also depends strongly on the nature of the
matrix. For example, it did not proceed in polypyr-
omellitimide and PVDFmatrices, but does run in PET,
bisphenol A polycarbonate and polyvinylchloride
matrices. The sensitivity of the process to matrix and
oxidant is a consequence of formation of a transition
state including the aniline molecule, the elementary
unit of PET and the HOCl (HOBr) molecule, which is
antecedent to the starting reaction (aniline oxidation).
The final PANI/PET composite had a conductivity
,1024–1026 S/cm and high transparency. Later, it
was shown [116,117] by means of conductivity
Table 1
Heat of fusion, melting temperature and degree of crystallinity of
Nylon 6 and PANI-HCl/Nylon 6 [113]
Heat of fusion,
DHf (J/g)
Melting point,
Tm (8C)
Degree of
crystallinity,
Xc (%)
Nylon 6 55.28 217.4 21.0
PANI-HCl/
Nylon 6
49.40 215.5 18.7
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1711
measurements, standard AFM, conducting AFM, DRS
and EPR techniques that the presence of PANI in a thin
(,1 mm) surface/subsurface layer strongly affects the
properties of the film composite (20 mm). Thus,
standard AFM topographical images proved that the
undoped form of the composite, as well as parent PET,
had a flat surface, whereas after doping its relief was
highly disturbed and exhibited ‘mountainous’ features
(Fig. 2). As one can see, HClO4 inducedmuch stronger
changes than did HCl (Fig. 2c and d), though the dc
conductivity of PET/PANI·HClO4 was reduced by a
factor of two (compare 9.1 £ 1026 S/cm with
1.8 £ 1025 S/cm). It is assumed that the reason for
such the marked change in the surface relief is due not
only to the appearance of charge carriers on PANI
macromolecules, but also to the distribution of
counter-anions compensating their charges in a con-
ducting surface layer of the composite. In particular, it
is accepted [117] that this effect was caused by a
rearrangement in the packing of the amorphous part of
PET and PANI, induced by Cl2 and ClO42 anion
penetration. In this interpretation, the larger ClO42
anions, when penetrating into the polymer matrix,
deformed it more strongly than Cl2 anions.
These transformations were characterized by
strong changes in the DRS spectra of the materials.
Specifically, the transition from the undoped to the
doped form one is accompanied by an increase of the
high frequency peak associated with the conductivity
of the clusters leading to the appearance in DRS
spectra of two low frequency relaxation processes,
connected with interfacial polarization phenomena.
Similar relaxation behavior is observed with compo-
sites doped by HClO4 and HCl acids. However, the
dielectric losses in the case of HClO4 are much higher
in the low frequency region, probably due to the
stronger deformation of the matrix.
One may infer from these data [117] that even
aside from the physical–chemical interactions among
PANI, the dopant and the matrix polymer, the size of
the dopant anion should affect the performances of the
PANI conducting composite. Specifically, it concerns
the dependence of the rate of release of the doping
acid on the dopant anion size from doped PANI [118]
and its composites when they contact with water, or
used under operating conditions [119]. Thus, Neoh
et al. [119] found PANI·HCl/Nylon 6 composite films
are readily converted to the base form due to a loss of
counter-ions (Cl2) when immersed in water. In
contrast, when using the larger dopant sulfosalicylic
acid, the composite films did not convert to the base
form, even after extended exposure to water or under
simulated weathering conditions.
Another diffusion-oxidation method [9] is aniline
(or other monomer) polymerization in polymer
matrixes impregnated with an oxidant that also allows
preparation of PANI (polypyrrole) conducting com-
posites, but this seems not to be very practical.
Specifically, it can be realized through exposing the
matrix polymer (e.g. poly(acrylamide)) impregnated
with an oxidizing agent to hydrochloric acid vapor,
and then to the monomer vapor [120] or solution
[100]. The conductivity of the resulting composites
reached 1025 S/cm. As one can see the main difficulty
here is the presence in the final composite of inorganic
products of the oxidant reactions, which can affect its
water resistance, mechanical or/and other properties.
2.2.3. Electrochemical polymerization of aniline in
a matrix
Although electrochemical polymerization of ani-
line in large-scale technologies is not practical, it can
be useful for small geometry systems (sensors,
microelectronics and optic devices, batteries, etc.),
due to such advantages as a strict control of PANI
properties produced at an electrode surface, the
possibility to avoid by-products of the process, etc.
[8,121]. Polymerization at an electrode (anode)
Fig. 2. 3D-AFM topographical images (TappingMode) of pure PET
and (a) the PANI/PET composite; (b) the undoped form; (c) the
PANI·HCl/PET form; (d) the PANI·HClO4/PET form [117].
Reproduced from Pud et al. by permission of J Mater Sci
2001;36(14):3355. q 2001 Kluwer Academic/Plenum Publishers,
New York, NY.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531712
surface coated by a non-conducting polymer film at
the aniline oxidation potential results in the formation
of a PANI/polymer composite [31,122]. The necess-
ary condition here is penetration (diffusion) of aniline,
solvent and electrolyte through the coating to its
interface with the anode [31], to create the electro-
chemical prerequisites to oxidize molecules of aniline
(in reality anilinium cations), and growing PANI
macromolecules. This condition can be realized in
two ways: (1) through pores and (2) by swelling the
polymer coating in the reaction medium (solution),
separately or in parallel, dependent on the coating
porosity and swellability. Under appropriate con-
dition, the polymerization starts in the interface
between the anode surface and the coating [31], and
the resultant PANI grows from this interface into the
coating bulk, forming a new electrically conducting
alloy film, as shown for different matrixes in the
polypyrrole case [27–30,123]. Whereas the polym-
erization process is directed from anode to cathode in
these electrochemical systems [124], the final com-
posite film will have a gradient of the conducting
polymer distribution in the insulating matrix. For
example, Wang et al. [125] found differences between
the solution and electrode composite film sides for a
composite with polypyrrole.
Some early works on the electrochemical polym-
erization of aniline in different polymer matrixes (PU,
PC, PMMA, poly-p-phenylene terephthalamide/di-
phenylether, etc.) were discussed in a review by
Anand et al. [31]. It was found that the conductivity of
the composites was close to that of pure PANI.
A template electrochemical polymerization of ani-
line in porous PVDF or sol–gel silica films covering
the anode surface produced micro- and nanocompo-
sites containing the conducting polymer with spectral
and electrochemical properties near to those of pure
PANI [126]. The surface morphology of these
composite films was described using the fractal
dimension concept [127]. Applying a similar polym-
erization technique with a cellulose acetate membrane
on a platinum electrode, das Neves and De Paoli [128]
produced PANI dispersed inside a microporous
membrane structure. It is appears that the photocurrent
response of the electrochemically synthesized PANI in
the pores of the membrane is enhanced in comparison
with a PANI composite film prepared by casting.
This suggests effects on the properties of the final
conducting composite from the PANI dispersity and
the interaction of the membrane host polymer with
PANI (compare with the chemical Section 2.2.2).
The effect of the matrix nature on the electro-
chemical polymerization of aniline and the properties
of the produced PANI composites was also shown by
Pud et al. [129] Specifically, they found that the
polymerization rate in the PA-12 matrix was higher
than in PVA, due to the stronger interaction of PANI
(aniline) with the PA-12 matrix. It was concluded by
spectral data that this process resulted in formation of
shorter chains PANI than for PANI formed in ‘free
conditions’ at a bare electrode. Moreover, it was
supposed that the interaction of aniline and PANI with
the matrix polymer hinders protonation of imine sites
in PANI. This can increase the response time of these
composites to different influences, particularly when
using these materials as sensor elements. However,
such an effect can be minimized by adjustment of the
electrochemical parameters of the aniline polymeriz-
ation [129] and varying the matrix. This is probably in
agreement with data of Bartlett and Simon [130] on the
electrocatalytic properties of electrochemically pro-
duced PANI/polyacrylic acid and PANI/poly(vinyl-
sulfonate) films at the anode for NADH (reduced
nicotinamide adenine-dinucleotide) oxidation. In com-
parison with films of PANI/poly(vinylsulfonate), the
amperometric responses of PANI/polyacrylic acid
were reduced by one third, the currents saturate at
lower NADH concentration, and the response was less
stable towards repeated measurements. At the same
time Andreev [131,132] reported that PANI films
produced electrochemically inside a Nafion film
covering Pt or glass carbon electrode mainly retained
its properties.
The effect of the composition of the reaction
mixture solution on the properties of electrically
conducting PANI/polyacrylonitrile (PAN) composite
films prepared by electrochemical polymerization of
aniline on the PAN-coated Pt working electrode in the
acetonitrile/water mixture solution was investigated
by Park and Park [133]. An acetonitrile (50%)/water
(50%) mixture was the optimal composition of
the solution in the preparation medium for the
dissociation of electrolyte (acid) and the transpor-
tation of aniline and electrolyte ions through PAN to
the working electrode. This suggested that the
optimum solution composition favored sufficient
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1713
swelling of the polymer film at the electrode to
enhance the rate of the aniline polymerization. The
maximum peak current was obtained with sulfuric
acid as an electrolyte. The electrical conductivity of
PANI/PAN composite film peeled from the Pt
electrode was around 1021 S/cm.
2.3. Polymer grafting to a PANI surface
As one of the stand-alone synthetic methods to
obtain PANI composites we may consider probably
the grafting of some polymers to a PANI surface.
Thus, Chen et al. [134] demonstrated chemical
modification of EB via its UV-induced surface graft
copolymerization with methoxy-poly(ethyleneglycol)
monomethacrylate macromonomer (molecular weight
,2000) in aqueous media. They showed that
modified PANI films doped by HClO4 were very
effective in reducing protein adsorption and platelet
adhesion. The authors [134] believe that these
materials greatly extend the potential applications of
PANI composites as biomaterials and blood compa-
tible materials.
Using thermally initiated graft copolymerization of
another acrylic monomer (acrylic acid), Chen et al.
[135] modified the surface of EB films or powders, to
create conditions for covalent immobilization of
enzyme invertase on the electroactive polymer
substrate. In their opinion, this method might provide
additional advantages over the conventional polymer
substrates for enzyme immobilization. Specifically,
the accompanying changes in the substrate redox
potential and conductivity after the enzymatic reac-
tion may give additional means for effective sensing
and detection of some enzymatic reactions.
The grafting method has opened also a possibility
to resolve the problem of the poor adhesion properties
of PANI. For example, Ma et al. [136] have developed
interfacial thermal graft copolymerization induced
laminations of EB/Polytetrafluoroethylene (PTFE) in
the presence of either acrylic acid or 1-vinylimidazole
monomers. Before the graft procedure the PTFE
surface was activated in argon plasma. The EB and
PTFE films were then lapped together in the presence
of a small quantity of the pure or aqueous monomer
solution and kept at 100–140 8C. This allowed a
maximum lap shear adhesion strength approaching
200 N/cm2 for the EB/PTFE interface laminated at
140 8C for 1 h in the presence of pure acrylic acid, and
with 40 s of argon plasma pretreatment time for the
PTFE surface.
Zhao et al. [137] developed the graft copolymer-
ization of vinylbenzylchloride (VBzCl) to PANI
base using UV- and heat-induced methods, which
resulted in the alkylation of the imine nitrogen of
PANI with VBzCl. The PANI conductive doped
state was obtained due to the formation of chloride
anions formed during the alkylation and acting as
the counter-anions to the N þ component. On the
other hand, the VBzCl polymerization via the vinyl
groups led to the formation of a hydrophobic layer
on the PANI surface. This layer acted as a barrier
preventing the undoping of the graft copolymerized
samples, which maintained their conductive state,
even when exposed to aqueous solutions with high
pH (,12) [137].
3. Blending methods
3.1. Solution blending
Together with the methods considered above, the
development of solution methods to process PANI is
based on the understanding of the fact that difficulties
in its processibility are related to its aromatic
structure, interchain hydrogen bonds and effective
charge delocalization in its structure [138]. These
difficulties have been overcome by approaches
imparting PANI dissolution in different solvents to
dissolve PANI and facilitate the preparation of PANI
conducting composites with polymers soluble in the
same solvents:
1. The synthesis of substituted polyanilines, which
are soluble in organic solvents, realized through
the introduction of alkyl [139,140], alkoxy [141]
and other substituents [14] on the monomer
benzene rings.
2. The introduction of sulfonic groups on PANI
benzene rings, to form water soluble sulfonated
self-acid-doped PANI (SPAN) [142,143] or highly
sulfonated SPAN [144]. Another kind of sulfo-
nated PANI can be produced through substitution
of hydrogen in imine sites of PANI, e.g. by
propanesulfonic acid (PAPSAN) [145].
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531714
3. Copolymerization of aniline with other monomers
to form soluble aniline copolymers [146].
4. Protonation (doping) of undoped PANI by
functionalized protonic acids (e.g. CSA, DBSA,
phosphoric acid diesters, etc.), generally denoted
as Hþ(M2–R), in which the counter-anion
(M2–R) bears a charge (e.g. sulfonic, substi-
tuted phosphoric, etc.) and an R– functional
group (e.g. alkyl substituted aromatics, long
alkyl chains, etc.), imparting dopant compatibil-
ity with non-polar or weakly polar organic
solvents [2–5]; Cao et al. [2] called this method
to dissolve PANI ‘counter-ion induced processi-
bility’.
5. The use of amide solvents such as NMP, in
which PANI base is soluble. This provided the
first serious success in solving the PANI
processibility problem [147].
The use of an acid–base interaction of the
PANI base with concentrated acids [148–150] or
solvents having a strong acidic function (e.g. hexa-
fluoro-2-propanol) [151].
3.1.1. Blends of substituted PANI
As frequently occurs, a gain in one aspect may be
accompanied by a loss in another. Thus, substituted
soluble polyanilines are characterized by decreased
conductivity in comparison with unsubstituted PANI
[14]. This can be explained by a disturbances of the
electronic delocalization in the polymer chains [152].
Nevertheless, the conductivity of PANI with some
non-bulky substituents is high enough for some
purposes. Thus, Dao et al. [153] reported a conduc-
tivity of 0.3 S/cm for POT and PMT upon doping with
HCl. Cazotti and De Paoli found POMA doped with
HCl exhibited even better electrical conductivity in
the range of 0.01–3 S/cm, dependent on the prep-
aration parameters [154]. Moreover, Raghunathan
et al. [155] found that an electron localization length
was much larger in poly(o-alkoxyanilines) compared
with corresponding poly(o-alkylanilines).These
potential POMA capabilities encourage investigators
to use alkoxyl substituents, mainly POMA, for their
solution blending.
Malmonge and Mattoso [156–158] used POMA
solubility to develop and study its film blends with
PVDF cast from blend solutions in dimethylacetamide
at various ratios of POMA-TSA/PVDF. The blend
composition had a great influence on the morphology
obtained. Specifically, at low content (5 wt%) of
POMA-TSA the morphology presented the growth of
fibrilles located preferentially in the boundaries of the
PVDF spherulites. On increasing the POMA-TSA to
10%, an interconnecting fibrillar-like morphology
was formed, and the spherulites characteristic of pure
PVDF could be hardly noted. For higher POMA-TSA
content, spherulites were not observed, and the
morphology consisted predominantly of intercon-
nected fibrils of diameter around 700 nm, spread
throughout the entire surface of the blend. Never-
theless, X-ray analysis confirmed the presence of the
b-crystalline phase characteristic of PVDF within the
blends, in addition to the presence of the POMA
component, which, at least for content below 25 wt%,
did not affect the b-PVDF structure. The growth of an
ordered structure with the main peak at ca. 2u ¼ 7:588could be observed for POMA content above 25 wt%.
Furthermore, the endothermic fusion peak assigned to
crystalline crystals of PVDF was observed for the
content up to 50–70%. These results suggest that the
crystalline structure of PVDF is affected upon POMA
addition in two forms: first, on the spherulites for
POMA content from 10 to 25%, and then on the
crystalline lamellae for POMA content in excess of
about 70% [156].
Mattoso and Malmonge [158] have also studied the
thermal behavior and electrical conductivity stability
of POMA-TSA/PVDF. They found that, unlike the
high thermal stability of PVDF up to 400 8C [159],
POMA-TSA weight losses proceed at much lower
temperatures, in a three-step process. The first step,
starting practically at room temperature and going up
to 130 8C, corresponds to the expulsion of imbibed
water from the polymer matrix. The second step, in
the range of about 220 8C up to 270 8C, is associatedwith dopant elimination and degradation reactions,
consistent with the boiling point of the dopant
(241.6 8C). Similar data were found by other teams
for PANI and its derivatives with different acid
dopants (HCl, H2SO4, H3PO4, HCOOH, TSA, etc.)
[160–164]. The third step, commencing at 270 8C, isassigned to degradation of a PANI chain, in agree-
ment with the literature [164,165]. It is encouraging
that the weight losses of POMA-TSA (e.g. 3.8% at
100 8C, 6.5% at 150 8C and 7.3% at 200 8C) were
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1715
diminished in the blend (0.7, 1.2 and 1.3%, respect-
ively). These losses were less than would be expected
for the 25 wt% content of POMA-TSA in the blend
(1, 1.6 and 1.8%, respectively) [158]. Moreover, the
onset temperature of these three weight losses usually
occurred at temperatures 5–10 8C higher than that for
pure POMA. For the POMA-TSA/PVDF blend, two
exothermic peaks at ,250 and 300 8C were observed
in the DSC analysis data. In accord with literature
reports [164,166], these are associated with degra-
dation reactions of the polymer chain and dopant
structures, such as cross-linking, loss of conjugation,
oxidation, decomposition and other reactions, includ-
ing a possible chemical reaction between the dopant
and the polymer. For the undoped blend the peak at
temperatures ,300 8C was not observed, probably
owing to the absence of dopant [158].
A low percolation threshold was observed with the
onset of conductivity at low POMA-TSA content (i.e.
5 wt%) [167]. The composite conductivity at ambient
conditions was about 1025 S/cm for POMA-TSA
content of 4.5 wt%. It was quite stable at temperatures
between 70 and 90 8C for the time scale studied
(500 h), with only a small decay during the first hours
of treatment, probably due to the elimination of the
residual solvent and/or water, which may contribute to
an increase in the charge carrier mobility, and
consequently in the conductivity [158]. Similar
changes for pure polyanilines were observed by
Javadi et al. [162] and Angelopoulos et al. [168],
who reported that small amounts of water lead to an
increase in the conductivity of polyanilines. They
explained this through a decrease in the apparent
separation of the metallic islands and/or the height of
the barrier between them, making tunneling more
favorable. On the other hand, a decrease in the
conductivity of the POMA-TSA/PVDF composite
films from 1023 to 1027 and 1029 S/cm for tempera-
tures of 130 and 150 8C, respectively, was attributedto dopant loss, degradation reactions, and structural
and morphological changes [158]. Specifically, treat-
ment at higher temperatures (130–150 8C for 500 h)
led to disappearance of the exothermic peaks in
the DSC spectra of POMA-TSA and composites,
indicating that degradation and dopant loss had
already taken place during the long treatment. These
transformations were confirmed by the fact that under
this treatment these samples became insoluble, and
remarkably less conductive. The insolubility indicated
the occurrence of a cross-linking processes at elevated
temperature.
Wilson et al. [169] prepared blend films of a
POMA-TSA/poly(epichlorohydrin-co-ethyleneoxide)
(Hydrin-Cw) rubber composite by casting from DMF
solution. The films at 9.1% (w/w) of POMA-TSA
content had an increase in conductivity by three orders
in compare with the parent rubber, without significant
changes in mechanical behavior. At higher PANI
content the composite Young’s modulus increased
nearly proportionally to the POMA-TSA content in
the mixture, reaching a maximum when the blend
contained 23.1% (w/w) of POMA-TSA. For this
blend, the elongation at break was about 300%.
POMA-TSA acts as reinforcing filler in the mixture,
making the rubber harder and less elastic. The
elongation at break decreased continuously with an
increase of the POMA-TSA content in the mixture,
and for the 50% (w/w) blend it was only 12%. Despite
this poor elasticity the films were flexible and self-
supporting. The electrical conductivity of polymer
blends was explained by the percolation theory, based
on the formation of a network of conductive material
in the insulating matrix. Specifically, Wilson et al.
[169] claimed that completely miscible mixtures are
not desirable because a conductive network will not
be created. The percolation threshold is defined as the
minimum amount of conductive filler which must be
added to an insulator matrix to cause the onset
of electrical conductivity. According to theoretical
studies, this occurs when the filler represent 16% (v/v)
in the mixture But a 10-fold increase in the electrical
conductivity was reported for 1.9% (w/w) of POMA-
TSA (dPOMA-TSA ¼ 1:07 g/cm3, drubber ¼ 1:32 g cm3,
thus 1.9% (w/w) ¼ 2.3% (v/v)). Therefore the perco-
lation threshold was at least 7 times smaller
than predicted by theory. The advantage of POMA-
TSA as a conductive component is the fact that
conductivity of its blends increases continuously. This
means that the level of conductivity can be modulated
by POMA-TSA loading, according to a desired
application [169].
Naturally, the properties and ease of preparation of
blends of alkoxy substituted PANI from a solution
depend on the solvent nature. Thus, Goncalves et al.
[170] investigated the suitability of different solvents
to prepare PU–POMA blend films by casting.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531716
Specifically, DMF, NMP andm-cresol were compared
for this purpose, with DMF selected as the most
suitable solvent for two reasons: suppressed deproto-
nation during the preparation of a predoped POMA
solution in DMF as compared with NMP, due to a
lower basicity of DMF as against NMP; convenience
in the use of DMF owing to its lower boiling point
(153 8C) than NMP (202 8C) or m-cresol (202 8C).Using PU–POMA solutions in DMF at different
weight ratios, Goncalves et al. [170] obtained flexible
conducting free standing films, which showed an
increase of conductivity from about 1026 to 1023 S/
cm as the POMA content in the blends changed from 5
to 65 wt%, respectively. The use of POMA predoped
in its powder form (from acid aqueous solutions) led
to POMA/PU blends with higher conductivity than
those in which POMA was doped in DMF solution,
independently of the dopant used. Films with POMA
doped in DMF solution were relatively fragile and
brittle for POMA content above about 40 wt%,
and rubbery but intractable for POMA content
below about 20%, for which the conductivity was
below 1026 S/cm. The films with predoped POMA
showed flexibility similar to films of pure PU. The
conductivity of the blend composition for POMA
predoped with p-toluene sulfonic acid was higher
than that of the blend predoped with trifluoroacetic
acid [170].
Paterno et al. [171–173] demonstrated polyelec-
trolyte complexes of poly(o-ethoxyaniline) (POEA)
with sulfonated lignin (SL) in a salt form. These
complexes could be obtained both in dilute solutions
[171,172], and as alternating POEA and SL self-
assembled layers, prepared through alternate adsorp-
tion of the components during 3 min immersion in
aqueous solutions [171–173]. The complexes demon-
strated some striking properties in compare with pure
POEA. Specifically, due to the charge screening effect
of anionic groups of SL, the degree of POEA doping
increased in aqueous solution, with a weakened pH
dependence: POEA in the complex remains doped
even at pH 7.0 but in the individual state POEA
becomes dedoped for pH .5.0.
Interesting data have been obtained for composites
of PANI with alkyl substituents. Thus, Anand et al.
[174] developed and studied soluble POT and PMT
blends with polyvinylchloride (PVC). At a previous
stage they synthesized POT and PMT in a salt form by
chemical oxidative polymerization, using HCl,
HNO3, H2SO4, H3PO4 and CH3COOH as acids. The
polytoluidines dedoped to their base forms were
soluble in THF, which is also a solvent for PVC. It
was found that POT and PMT bases produced as salts
of HNO3 were the most soluble among other bases;
the authors did not discuss reasons for this phenom-
enon [174]. These bases were chosen for solution
blending. Blend solutions 2% (w/v) were precipitated
in petroleum ether (non-solvent), followed by drying
and doping with HNO3. TGA–DTA and DSC
measurements showed that the thermal and oxidative
stability of POT-HNO3/PVC and PMT-HNO3/PVC
blends (powders) were much better than those of
individual polytoluidines. However, conductivity of
the blends was vice-versa. Specifically, pure POT-
HNO3 had conductivity 1.7 £ 1023, which lowered in
its blends to ,1026 S/cm at its 50 wt% content. At
the same time, the dielectric constant ð10rÞ and
dissipation factor ðtan dÞ of the blends were higher
than those of PVC due to the presence of the
conducting polymer in the blend, and increased with
its content. However, the highest dielectric constant
obtained for POT(90)-HNO3/PVC(10) blend was
more than two orders of magnitude smaller than that
of pure POT-HNO3. It is interesting that for similar
blends in a PS instead of a PVC matrix, the composite
conductivity was better. Thus, with PS content up to
30 wt% there was no significant drop in conductivity
in comparison with pure PMT or POT [175]. For a
solution blended in formic acid composites of PMMA
with PMT or POT doped by formic acid Anand et al.
[176] confirmed that the thermal stability of these
blends was higher than that of the pure salts, as
reported for POT-HNO3/PVC and PMT-HNO3/PVC
blends [174].
Ahmed et al. [177] obtained better conductivity
when used picric acid as dopant for POT, and
produced its composite with ABS through solution
blending in m-cresol. A remarkable low percolation
threshold (,3 wt%) was demonstrated for this
composite, similar to that of PANI-picric acid/ABS
blends (,4 wt%) [178]. The conductivity of films
with POT-picric acid content above 10 wt% was
about 1 S/cm [177]. This nice result was probably
due to the use of m-cresol as the solvent, which,
according to MacDiarmid and Epstein [179], affected
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1717
the polyaniline chains conformation to facilitate
charge transfer.
Sevil et al. [180] demonstrated that chlorine
substituted PANI had enhanced solubility in compari-
son with pure PANI. Specifically, they prepared 2-
chloro-polyaniline (2-Cl-PANI) in its non-conducting
EB form, and dissolved it with PVC in THF for
casting into thin composite films. The conductivity of
these films increased by more than four orders of
magnitude (from 1026 to 1022 S/cm) when they were
exposed to UV, g-rays and e-beams. This was
attributed to the subsequent doping of 2-Cl-PANI
with HCl released due to PVC the dehydrochlorina-
tion under radiation [180]. This phenomenon, which
can be anticipated for composites with variants of
PANI and a PVC matrix (or other halogenated
polymers with the ability to eliminate hydrohalogen)
may be likened to an internal composite self-doping.
On the other hand, the behavior also means that
in such composites PANI can act as a trap for HCl,
and hence it may be considered as some sort of a PVC
stabilizer.
Chen and Hwang [181] prepared PVA blends with
water-soluble self-acid-doped conducting polyani-
lines, specifically with sulfonic acid ring-substituted
polyaniline (SPAN) and poly(aniline-co-N-propane-
sulfonic acid aniline) (PAPSAH). They supposed that
the strong interaction of these polyanilines with PVA
through hydrogen bonding between hydroxyl groups
(of PVA) and positively charged amine and imine sites
(of SPAN and PAPSAH) led to a decrease in hydrogen
bonding amongPVAchains and to a partialmiscibility.
When the PVA content was higher than 70 wt%,
interconnected regions of PVA-rich phase and of
SPAN-rich phase were formed such that the dilution
effect of PVA on the conductivity was not large [181].
These observations suggest applicability of these
composites at different loadings of the conducting
phase in water systems. Specifically, due to the
composite water swelling capacity and electrocatalytic
properties inherent to conducting polymers [10] they
can find possibly an interesting application in sensor
devices, as shown for an PANI–poly(vinylsulfonate)
composite by Bartlett and Wallace [182].
3.1.2. Blends of soluble aniline copolymers
In this part we consider PANI block or grafted
copolymers with blocks of more conventional
dielectric polymers. Although these materials could
have been discussed above, they are considered
separately as their solubility opens an additional
method to produce blends with soluble common
polymers. Unexpectedly, some of the copolymers
displayed conductivities as high as the best samples
discussed in the preceding, up to a few units S/cm,
indicating an adequate length of conjugated PANI
blocks in the macromolecular chain. This supposition
is in accord with data of Lu et al. [183] for PANI
oligomer-phenyl capped octaaniline, for which the
conductivity was the same order of magnitude as
higher molecular weight PANI.
Generally, this method is based on the ability of
aniline or imine units of EB or LEB to interact with
reactive end groups of dielectric polymers. Li et al.
[184] were probably the first to develop routes for
PANI solubilization by the synthesis of soluble aniline
copolymers. They synthesized A–B–A block copo-
lymer, with segment A the PANI block and segment B
a poly(ethyleneglycol) with a –C6H4–NH2 end
group. The conducting block copolymer was formed
by the slow addition of aniline to a solution of
polymer and oxidant, extending their earlier studies
on the synthesis of graft copolymers of PANI [185]. In
particular, the PANI chain grew from the –NH2
pendant group of a carbochain polymer with a flexible
saturated chain, such as poly( p-aminostyrene) or
poly(vinylamine). In another route, the reaction of
amine alkylation was used to produce graft copoly-
mers of PANI. Thus, the graft copolymer was
obtained by refluxing EB and poly(epichlorohydrin)
dissolved in cyclohexanone in the presence of sodium
hydroxide. Neutral copolymers obtained by
titration with aqueous NH4OH solution were soluble
in DMF, DMSO and THF and slightly soluble in
CHCl3 and CH3OH, while the protonated copolymers
were much less soluble in these solvents. The
conductivity of these copolymers was in the range
of 1024–1 S/cm.
The synthesis of polyaniline ABA triblock copo-
lymers soluble in organic solvents was also carried out
by Kinlen et al. [186], with a diamine (–NH–C6H4–
NH2) terminated polymer as the B block and PANI as
the A block. The authors [186] proceeded from the
premise that the diamine moiety was more easily
oxidized than aniline. Accordingly, they assumed that
the first step in the reaction was the formation of
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531718
a short lived amine radical cation, which nucleophi-
lically attacked aniline at the para position. Copoly-
mers were synthesized starting with diamine polymers
of polyethyleneoxide, polypropylene oxide, polydi-
methylsiloxane and polyacrylonitrile-co-butadiene
with molecular weight ðMwÞ in the range of 400–
5000. These polymers were added to aniline–
dinonilnaphthalenesulfonic acid emulsion prior to
ammonium peroxydisulfate addition. Finally, this
method resulted in soluble ABA triblock copolymers
with molecular weights ðMwÞ from 30,000 to 140,000.
Yamaguchi et al. [187] used a typical amine-
epoxide reaction when treating LEB with phenylgly-
cidylether (PGE). That caused a ring-opening
polymerization of PGE, to result in a graft copolymer
(LEB-g-PGE) that was soluble in acetone and
chloroform, which are poor solvents for LEB. The
LEB-g-PGE/LiBF4 composite film was obtained
through the evaporation of a dimethylformamide
solution containing LEB-g-PGE and LiBF4. This
film showed lithium ionic conductivity of
1.0 £ 1026 S/cm at 295 K. The use of LEB opens
interesting synthetic possibilities in producing corre-
sponding soluble copolymers or their blends. How-
ever, owing to its oxidative instability, this method
should include additional measures to stabilize the
polymer during synthesis and operating conditions.
A water-soluble self-doped poly(aniline-co-2-
acrylamido-2-methyl-1-propanesulfonic acid (PAM-
PANI) copolymer with good conductivity was
prepared by Yin and Ruckenstein [188]. N-(4-
Anilinophenyl)methacrylamide (APMA) was syn-
thesized via the catalytic aminolysis reaction (from
p-aminodiphenylamine and methylmethacrylate), and
poly(AMP-co-APMA) (AMP ¼ 2-acrylamido-2-
methyl-1-propanesulfonic acid) was obtained through
a surfactant-free emulsion polymerization in water for
use in a graft copolymerization of aniline onto the
aminodiphenylamine pendant groups of poly(AMP-
co-APMA). The PAMPANI film cast from
water solution gave the remarkable conductivity of
about 4 S/cm.
3.1.3. Blends prepared due to counterion-induced
solubility of PANI
As mentioned above, the discovery of a processing
route for the conductive form of the PANI-emeraldine
salt of functionalized sulfonic acids by Cao et al. [2]
marked a significant advance. These acids induce
solubility of doped PANI in non-polar or weakly polar
solvents. Specifically, films cast of PANI-CSA (at
their 2:1 molar ratio) m-cresol solution had a
conductivity of ,400 S/cm. These secondary doping
phenomena were attributed to an expanded coil-like
conformation, which was proven by viscosity studies
[179,189]. Ikkala et al. [190] believe that this
conformation resulted in supramolecular structures
due to the combination of three specific simultaneous
interactions: first, the sulfonic acid is bonded to PANI
through proton transfer; second, the hydroxyl group of
m-cresol forms a hydrogen bond with the carbonyl
group of CSA; and third, the phenyl groups of m-
cresol and PANI stack, yielding enhanced mutual
dispersion forces. They also have shown that such the
specific interactions are allowed by the molecular
dimensions and by steric details simultaneously, thus
providing the requirement for what was called
‘molecular recognition’. Such interactions promote a
more extended conformation of the PANI chains,
which leads to the improvement in solubility and
conductivity [179]. The improvement in conductivity
persists even after eliminating the residual m-cresol
from the cast PANI films. This suggests that the chain
structure determined by the solvent and formation
conditions is maintained in the solid PANI. Indeed,
as Kugler et al. [191] have observed, the submono-
layer coverage of EB/CSA spin-coated from a
chloroform solution contain geometrically shaped
crystalline islands, whose internal structure was
attributed to the presence of compact coils. Upon
secondary doping using m-cresol vapor, the crystal-
linity was lost.
Considerable interest has been focused on this
processing route, including the possibility of stretch
aligning of the PANI doped fibers and films,
resulting in increased conductivity in the stretch
direction up to 1000 S/cm [191,192]. The founders
of this process demonstrated that a complex of
PANI with functionalized sulfonic acids could be
processed in blends with common insulating
polymers, such as PMMA, PC, Nylon 4,6 and
Nylon 12, polyvinylacetate, polyvinylbutyral, ABS
by preparation of their joint solution, followed by
film casting [2,5] to obtain blend materials with
interesting characteristics. Specifically, PANI-CSA/
PMMA displayed probably the most unique
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1719
conductivity, transparency and other properties.
Thus, Yoon et al. [193] found an extremely low
percolation threshold ðfp < 0:003Þ when investi-
gating the transport properties of PANI-CSA/
PMMA cast from m-cresol. The electrical conduc-
tivity of these blends followed the Mott–Deutscher
model for variable-range hopping on fractal net-
work, sðTÞ , exp½2ðT0=TÞg�: The g coefficient
increased from 1/4 in pure PANI-CSA (inducing
variable-range hopping among exponentially loca-
lized states) to 2/3 as the PANI-CSA concentration
was reduced to fp [194]. Moreover, the characteristic
metallic properties of pure PANI-CSA (positive
temperature coefficient of resistivity at high tem-
perature, linear temperature dependence of the
thermoelectric power, and frequency independent
ac conductivity) were retained in PANI-CSA/
PMMA blends down to fp [193,195]. Optical
quality, transparent conductivity films of PANI-
CSA/PMMA combined low surface resistance with
excellent transparency [196,197]. For example, films
were prepared with surface resistance less than
100 V/A and transmittance of ,70% between 475
and 675 nm. Their transmission electron micro-
graphs revealed the formation of an interpenetrating
network of fibrillar, crystalline PANI within the
PMMA matrix [198]. In the dilute regime, the PANI
morphology was a tenuous interconnected fibrillar
network, with a characteristic cross-sectional fibril
dimensions of a few tens of nanometers. Fraysse
et al. [199] showed that existence of the inter-
connected network affected also thermomechanical
properties of PANI-CSA/PMMA and PANI-
DEHEPSA/PMMA composites prepared from m-
cresol and dichloroacetic acid, respectively. Thus,
whereas the matrix underwent an irreversible flow
slightly above its glass-rubber transition tempera-
ture, blends with PANI-CSA mass fraction as low as
1 wt% showed a well-defined rubber plateau, with a
tensile modulus in the MPa range for temperature in
the 400–500 K range (Fig. 3). The authors [199]
interpreted the results as an example of ‘mechanical
percolation’ and concluded that the mechanical
percolation threshold (0.5–1 wt%) of the blend
was significantly higher than the electrical one
(0.04–0.07 wt%). However, they emphasized that
this effect requires more accurate experiments to be
well characterized, and this may be important for
understanding the real effect of PANI on the blend
properties.
Jousseaume et al. [200–203] investigated the
evolution of transport properties with temperature for
blends of PANI doped by CSA or DiOHPwith PS, cast
from m-cresol solutions. Using a fluctuation induced
tunneling model, they explained the electrical conduc-
tivity variations of the blends in the temperature range
between 77 and 300 K by a hopping mechanism
between conducting clusters separated by thin insulat-
ing barriers. Above the percolation threshold
(.1 wt%) thermal aging of the blends led to an
expansion of the insulating barriers, to the detriment of
the cluster size. It was showed that the aging kinetics of
PANI-DiOHP films was faster than that of PANI-CSA
films, but that these pure PANI-DiOHP and PANI-
CSA films and corresponding blends exhibited similar
kinetics of thermal aging at the same temperature
[200–202]. For condition near room temperature
Jousseaume et al. [203] confirmed the ‘metallic’-type
behavior of these blends and pure doped PANIs above
the percolation threshold. However, they observed an
irreversible degradation of PANI-DiOHP/PS and the
blends of PANI-CSA at temperatures near 450 and
500 K, respectively, that resulted in a large decrease of
their conductivity. It is interesting that the same
Fig. 3. Thermal dependence of the storage modulus (logarithmic
scale) for PANI-CSA/PMMA blends with PANI content between 0
and 100% [199]. Reprinted with permission from Fraysse et al.
Macromolecules, 2001;34(23):8143. q 2001 American Chemical
Society.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531720
degradation temperature was observed for pure doped
PANIs, showing that the conducting networks of the
doped PANI in these blends were stable up to the
degradation temperatures PANI [203].
Naturally, the cast films of doped PANI and its
blends can retain solvent. This is important, especially
in the case of high boiling solvents, which are very
difficult to completely remove. In turn, this may effect
several material properties. For example, Jousseaume
et al. [203] revealed a decrease of the conductivity by
electrical conductivity measurements during heating–
cooling cycles, as the residual solvent (m-cresol) and
moisture evaporation of. This phenomenon was
explained by the existence of a frontier sensitive to
the solvent at the periphery of conducting clusters.
Specifically, for PANI-DiOHP/PS films, the tempera-
ture dependence of conductivity before and after the
partial evaporation of the solvent was well described
by a model of a tunnel effect limited by the charging
energy of conducting clusters.
Obviously, these specific solvent-conducting clus-
ter interactions are sensitive to the dopant, and can
lead to differences of conductivity of doped PANI
[204] or its blends [205]. Kuramoto and Teramae
[205] showed that PANI-DBSA/PMMA composites
prepared from m-cresol solutions at PANI content of
,0.4 and 10 wt% had conductivity ,1025 and
,3 £ 1023 S/cm, respectively, that was much lower
than for PANI-CSA/PMMA [193] with corresponding
conductivities ,1022 and 0.2 S/cm, respectively.
It should be noted that these interactions do not
exhaust all solvent effects on the properties of the
PANI blends prepared by solution casting. As in the
case of coatings prepared from solvent based paints, it
seems that the properties and the quality of the cast
films of PANI and its blends depends on the physical
and chemical characteristics of the solvent, on the
complex of interactions involving all the solution
components, the surface of the substrate support, and
the preparation conditions. For example, Valenciano
et al. [206] prepared blends of PANI doped by CSA
with UHMW-PE in a solvent mixture of m-cresol and
decalin. They showed that the preparation conditions
were very critical to obtain a high quality film cast
from the mixture. In particular, a change in the solvent
mixture proportions or in the dopant could result in a
non-cohesive or very brittle film. For example, to
obtain the desired composition, UHMW-PE was
dissolved in decalin (,5 wt%) and added to a
PANI-CSA0.5 solution (,1 wt%) in m-cresol, keep-
ing the m-cresol to decalin ratio at 1:2.4 (v/v). This
produced homogeneous and flexible films, with a low
percolation threshold (,1 wt%), and electrical con-
ductivity of 1026 and 0.01 S/cm, for blends contain-
ing 1 and 5 wt% of PANI, respectively. The tensile
strength of the UHMW-PE film (3.3 MPa) could be
maintained in the blend up to the 10 wt% PANI, after
which it dropped drastically. The elongation at break
of UHMW-PE, which was usually above 400%,
significantly decreased with the addition of PANI.
To explain the mechanical properties of the blends,
the authors suggested a phase separation due to a
saturation of PANI concentration in the blend,
confirmed by electrical conductivity changes and
preliminary SEM studies. The observed decreases in
the heat of fusion and melting temperature were
consistent with some degree of miscibility of PANI-
CSA with the UHMW-PE [206]. These changes in the
thermal features may be explained by the effects of
PANI on the blend morphology during the solvent
evaporation. Indeed, Zhang et al. [207] have observed
that the blend crystallinity decreased in comparison
with the parent polymers for PANI-CSA blends with
different polyamides (PA-66, PA-11, PA-1010), cast
from a m-cresol–chloroform (50/50, v/v) mixture,
with a corresponding decrease of the blend melting
temperatures and entropies.
In turn, these observations suggest that rigid rod-
like PANI macromolecules can hinder the packing a
matrix polymer into crystallites on the formation of
the solid from the blend solution or melt. Moreover,
the crystallization process can also be effected by the
positive charge of PANI doped macromolecules and
their molecular mass, and by the size and nature of the
doping agent. Thus, the extent, or absence, of doping
of PANI may alter the degree of the crystallinity in the
blend and the form and size of crystallites (compare
with Ref. [116]), producing blends differing by their
transport, mechanical and other properties. All these
effects can be amplified by PANI loading as well [207,
208]. For example, a significant drop in crystallinity
with increasing PANI fraction from 0 to 9 wt% was
reported by Zhang et al. [207] for PANI-CSA/PA
blends. In the case of blend fibers of PANI-DBSA and
UHMW-PE prepared by Andreatta and Smith [208]
through solution blending in decalin with various
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1721
ratios of PANI to UHMW-PE, the modulus and the
tenacity of the fibers ranged from 40 to 0.5 GPa, and
from 2 to 0.02 GPa, respectively. Conductivity was
3 £ 1024 S/cm for blends containing 5 wt% of PANI.
Leyva et al. [209] demonstrated an interesting and
surprising difference between rubbery blends of SBS
triblock copolymer and PANI-DBSA produced by
solution blending under magnetic stirring or ultra-
sonic vibration. Specifically, they found that blends
prepared in solution by magnetic stirring displayed a
higher conductivity than those obtained by sonication.
Basing on optical microscopy data the authors
deduced that the difference is connected with the
fact that sonication led to the formation of very small,
conducting particles well distributed inside the
matrix, while the magnetic stirring method gave
larger of PANI-DBSA particles. As a result the
sonicated system gave blends with higher percolation
threshold (3.8 wt%) than that for the magnetic stirred
system (2.2 wt%). This explanation contradicts the
usual guideline that at the same loadings, the higher
the dispersity of the conducting particles, the better
the formation of conducting pathways. However, we
may consider the results [209] as an important
indication on the necessity to control stirring con-
ditions to obtain reproducible results for PANI blends
obtained by solution blending.
As demonstrated above for polymerization systems
(Section 2.2.2), PANI blend properties should depend
on the ability of the matrix polymer to interact with
dopant and PANI. For example, Kim and Levon [210]
observed a homogeneous smectic liquid-crystalline
structures at PODA content below 10 wt% in a ternary
blend film cast from a chloroform solution of a
PANI(DBSA)4 complex with comb-shaped poly(oc-
tadecylacrylate) (PODA). This mesophase formation
was caused by the interaction between DBSA with
long alkyl chains of PODA and by hydrogen bonding
between the PANI complex and PODA. The effects of
component interactions were interestingly displayed
in the range of low PODA concentrations. Specifi-
cally, conductivity of the PANI(DBSA)4 complex
abruptly decreased from 9.9 £ 1022 to 1.3 £ 1023 S/
cm upon addition of 5 wt% PODA. Kim and Levon
[210] supposed that this was caused by interaction of
the PODA carbonyl groups with nitrogen cations
of the PANI complex, and indicated a localization of
electrons in short p-electron segments of PANI
complex chains. With increasing PODA in the ternary
blend, this interaction is reduced because of phase
separation between these components, and the
conductivity increased somewhat, e.g. to
3.8 £ 1022 S/cm for 10 wt% PODA content. The
conductivity did not change much at higher concen-
trations of PODA, and for the PANI(DBSA)4/PODA
30/70 blend, the conductivity was 0.02 S/cm. Among
factors affecting the properties of these ternary blends,
the authors [210] considered also hydrogen-bonding
interaction of the PANI complex with PODA and a
weak interaction of the methylene units of DBSA and
PODA.
It is known that PANI doped with a binary mixture
of sulfonic acids possesses peculiar thermostability,
conductivity and other characteristic features as
compared to the polymer doped separately by sulfonic
acids such as DBSA, TSA or naphtalenedisulfonic
acid [80]. Koul et al. [211] have shown enhanced
electrical and optical properties, along with higher
solubility in all common organic solvents, for PANI
doped with a mixture of DBSA/TSA (1:1). Using this
double doped PANI, they prepared composite films
with ABS by casting from the chloroform solution.
The surface resistance of these composites changed
from 300 MV/A to 1.302 kV/A, dependent on the
PANI doped content and the method of mixing the
system components.
As follows from the preceding discussion on PANI
composites obtained through aniline matrix polym-
erization, the ability of PANI to form hydrogen bonds
can affect properties of the final material. Naturally,
this is more intrinsic to polymers with polar groups in
their main or side chains than to less polar polymers.
This dependence provides a means to affect to some
extent the miscibility, mechanical, thermal, and
electrical properties of the conducting polyblends
through a change of the functional composition of the
matrix polymer. Various methods to enhance the
properties of immiscible blends include the use of
precursors, compatibilizers such as block copolymers,
or ionic polymer groups [212]. The last was used by
Ho et al. [213] when making a rubbery-like conduct-
ing polymer blend of thermoplastic PU (synthesized
from polytetramethyleneoxide and 4,4-methylene-
bis(phenylisocyanate)) with PANI-DBSA by
mixing in chloroform. The sulfonic chain extender
(2,5-diaminobenzenesulfonic acid) of PU allowed
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531722
additional hydrogen bonding with PANI-DBSA. As a
consequence, the blend composition led to variation
of the glass transition points, different degrees of
miscibility and a tensile strength of the modified PU
blends that increased with the incorporation of PANI-
DBSA.
A similar approach of introducing ionic groups,
such as a sulfonic moiety, to the insulating polymer
PC to enhance its coulombic interaction with PANI
phase in the composite was used by Lee et al. [214].
They reported on the preparation of conductive
flexible composites of PANI and sulfonated PC
(SPC), with improved compatibility of the com-
ponents. Thus, they found that the electrical conduc-
tivity of PANI-DBSA/SPC composites obtained from
chloroform solutions was larger by factor approxi-
mately 2–5 than that of PANI-DBSA/PC, and
increased to 7.5 S/cm at 25 wt% of PANI-DBSA
content. The percolation threshold of the PANI-
DBSA/SPC composite was about 15 wt%. The
conductivity difference suggested that the PANI-
DBSA complex might be distributed differently in
the two matrices, perhaps due to the effect of
sulfonation. However, in spite of this, the tensile
strengths of the two composites were almost the same.
It is interesting to note that the electrical conductivity
of PANI-DBSA/SPC composite was larger than that
of PANI-CSA/SPC.
Some differences of the conducting phase distri-
bution and properties were also observed for blends
of Nylon 6 or Nylon 12 (having the same polar
amide groups, but differing in the number of the non-
polar –CH2– units in their chains) with PANI doped
by CSA, DBSA, or MSA, cast from hexafluoro-2-
propanol [215–217]. Thus, Hopkins et al. [215]
analyzed the morphology of the conducting salt
component by small-angle neutron scattering data,
and analyzed this by two standard models for two-
phase systems: Debye–Bueche (D–B) and inverse
power law (IPL). At 3 vol% PANI-CSA0.5 concen-
tration the D-B model suggested salt domains with
characteristic lengths of 22 nm for the Nylon 12
blend. However, this differed from the blend with
Nylon 6, for which the IPL model indicated fractal
geometry. With increased content of the doped PANI,
modified structures were observed with both Nylon
blends [215]. This agrees with the finding that
significant molecular mixing is absent for mixtures
of Nylon 6 with deuterated PANI (D-PANI) [215].
Specifically, in the case of the lowest concentration of
D-PANI/CSA there was an indication of mass fractal
structure, but this was not found at higher concen-
trations. The results showed that blends with the
smaller and more polar dopants CSA and MSA
behaved similarly, but differently than either D-PANI/
DBSA blends or the D-PANI-emeraldine base. X-ray
scattering demonstrated the presence of Nylon 6
lamellae and residual peaks attributable to the pure
components [216]. Using differential scanning calori-
metry of PANI blends with Nylon 6 and 12 (dried at
110 8C), Basheer et al. [217] found that Tg was nearly
independent of the PANI and the sulfonic acid dopant
content, indicating a phase-separated morphology of
the blends. However, according to electron
microscopy data, the PANI domain size depended
upon the functionalized acid dopant, and that can
affect the blend conductivity. The decrease in the
melting transition temperatures of Nylon 6 and the
associated enthalpies with the blend composition was
attributed to the formation of the free acid dopant and
decomposition products of EB, which interacted with
the Nylon crystal content during thermal analysis
[217]. Hopkins and Reynolds [218] published inter-
esting data on the effect of the crystalline structure of
the matrix polymer on the formation of electrically
conducting networks in conducting blends. Specifi-
cally, blends of PANI-CSA with the crystalline and
amorphous polyamides Nylon 6 and Selar, respect-
ively, containing increasing PANI-CSA content,
showed conductivity with the rise more rapid for the
crystalline polymer. The authors concluded that this
faster conductivity rise stemmed from more devel-
oped conductive pathways in the crystalline host
blends, due primarily to the crystallization driven
exclusion of the conducting polymer into the inter-
spherultic region, as seen by transmission electron
microscopy. A conductivity of 1 S/cm was found for
PANI-CSA/Nylon 6 blends with 10 wt% PANI-CSA,
approximately 10 times higher than the conductivity
observed with the amorphous Selar matrix polymer.
Similar data were obtained for poly(3,4-ethylenediox-
ythiophene)/PSS blended with either PEO and PVA
[218]. These data [215–218] suggest that blends
containing an identical matrix polymer and equal
doped PANI loadings, but with crystallites differing in
size and quantity (degree of crystallinity) can have
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1723
different conductivity properties. Obviously, this is
determined by the distribution and orientation of
conducting PANI clusters among the crystallites in
the amorphous phase of the matrix polymer.
The importance of physical–chemical interaction
of the matrix polymer and doped PANI for conducting
blend properties was also demonstrated byWang et al.
[219] for PANI/PEO blends cast from aqueous
solution. They used an acidic phosphate ester dopant
prepared through reaction of POCl3 with poly(ethy-
leneglycol)monomethylether (PEGME, Mw ¼ 350).
The DSC curves of the blends with different doped
PANI loadings showed a shift of the single endother-
mic peak (at 67 8C in pure PEO) corresponding to a
suppressed melting temperature for the PEO crystal-
lites. This effect was explained by compatibilization
of the rigid conjugated polymer with the matrix
polymer, achieved due to the ability of the ester
dopant to form hydrogen bonds with PEO, reducing
the interfacial energy of the two incompatible blend
components [219]. This phenomenon may be con-
sidered to be a kind of plasticizing effect caused by the
long poly(ethyleneglycol) tail of the dopant. This
accords with the description of Geng et al. [220] for a
similar water soluble blend of poly(ethyleneglycol)
and poly(N-vinylpyrrolidone) with PANI doped by
phosphonic acid containing hydrophilic PEGME
ðMw ¼ 550Þ tails. Wang et al. [219] found that the
blends had an electrical conductivity percolation
threshold as low as 1.83 wt% PANI. Based on
conductivity and morphological studies, they con-
sidered the blend structure to be a three-phase system,
consisting of a crystalline phase of PEO, an
amorphous phase, and conducting PANI phase,
dispersed in the amorphous phase, leading to the
low percolation threshold by the double percolation
model [219]. In this work, an interesting negative
effect of the molecular weight of the matrix polymer
(PEO) on the blend conductivity was discovered.
Specifically, the conductivity dropped by two order of
magnitude as the molecular weight of PEO increased
from 20,000 to 5,000,000, for the same PANI loading
(3.19 wt%). The authors explained this by difficulties
in the PANI chain movement in the matrix polymer
with the higher molecular weight, hindering assem-
bling of a conductive pathway during transition of the
blend from solution to the solid state [219]. Differ-
ences in the distribution of the conducting phase in
the amorphous phase, whose condition and volume
fraction may change with molecular weight of PEO,
may also play a role. Unfortunately, comparison of the
crystallinity was not done for blends of PEO having
different molecular weights.
It naturally follows from the above that protonated
functionalized acids enhancing the solubility of PANI
in different solvents and/or compatibility with some
polymers, may be considered as potential plasticizers
of PANI and even some of its blends. Indeed, these
abilities appear to be due partly to incorporation of
dopant anions and molecules (in the case of the dopant
surplus) among the rigid rod-like PANI macromol-
ecules. This results in an increase of the intermole-
cular distance, a corresponding decrease of the
intermolecular interaction, and the PANI plasticiza-
tion. Specifically, Pron et al. [4,221–223] found that
protonating agents like phosphoric acid aliphatic
diesters induced solution processibility that allowed
formation from solutions of highly conducting PANI
blends with PS, ABS or PMMA at very small fraction
of the doped PANI. These neat diesters protonated EB
under mechanical mixing that resulted in a heavily
plasticized mixture which could be hot-pressed into
conducting, freestanding films. This ability would
make them a good choice to develop and produce
PANI blends and composites on an industrial scale.
However, at higher temperatures (above 140 8C),partial degradation of the PANI-phosphoric acid
diester complex occurs, leading to a decreased
conductivity, and a simultaneous increase of Young’s
modulus and tensile strength.
Using these diesters as plastisizing PANI dopants,
as long ago as 1993 Pron et al. [4] showed the
possibility of formation of conducting PANI blends
with DOP plasticized PVC with excellent mechanical
properties. Later, Pron et al. [224–226] demonstrated
this approach for highly transparent and conductive
PANI/cellulose acetate (CA) composite films cast
from m-cresol solutions. They compared properties of
the blends, either without plasticizer, or with common
plasticizers (dimethylphthalate, diethylphthalate and
triphenylphosphate). They tested different groups of
protonating agents: sulfonic acids, phenylphosphonic
acid, aliphatic (dibutyl and dioctyl) diesters of
phosphoric acid, aromatic (diphenyl, di-p-cresyl, and
di-m-cresyl) diesters of phosphoric acid. The compo-
sites of the doped PANI with unplasticized CA
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531724
demonstrated poorer mechanical properties and
higher percolation threshold (e.g. ,4 wt% in the
case of PANI-CSA) than those of the composites with
plasticized CA. The addition of plasticizers not only
improved the flexibility of the composite films, but
also significantly lowered the percolation threshold
(for the PANI-CSA/CA composite to fp ¼ 0:84 wt%,
for other blends to values below 0.5 wt% [224–226].
3.1.4. Preparation of PANI blends from solutions in
concentrated acids
Andreatta et al. [227] found that PANI readily and
completely dissolved at room temperature up to
20 wt% concentrations in concentrated (97 wt%)
sulfuric acid to yield homogeneous, viscous solutions
of a purple black color. They found no appreciable
degradation on repeated dissolution of PANI in
sulfuric acid and a precipitation in water or methanol.
This finding was in contrast to results generally
obtained with PANI film cast from NMP, which in
most cases could not be redissolved in NMP or in
sulfuric acid, indicating cross-linking of the PANI
macromolecules. These data show to the possibility of
forming composites of PANI with other polymers
soluble and stable in the same strong acids.
Specifically, composites of PANI with poorly
processible polymers may be produced by this
method. These composites usually contain PANI
doped by acid, which also serves as the solvent. For
example, Andreatta et al. [228] produced electrically
conductive blend fibers from a dilute isotropic
solution of PANI and PPD-T in 98 wt% sulfuric
acid. However, fiber produced with high PANI
content (25 wt%) had 233 MPa tensile strength,
which was not high enough for a wide range of
applications. Later, Hsu et al. [229] prepared a
composite fiber of PPD and the emeraldine salt of
PANI with better mechanical properties by mixing EB
polymer in PPD-T/H2SO4 spin dope, and extruding it
into green color fibers, containing only 1.0 wt% of the
emeraldine salt, with typical a diameter of approxi-
mately 25 mm. The fibers had an initial module of
62 GPa and a tenacity 2.8 GPa, compared with 76 and
3 GPa, respectively, for PPD-T fiber.
High strength and high modulus electrically
conducting PANI composite fibers were also prepared
by Hsu et al. [230] from air-gap spinning of lyotropic
PANI/PPD-T sulfuric acid solutions. The modulus
and tenacity of the composite fibers were in the range
of 28.6 and 1.7 GPa, respectively, for much higher
[229] PANI loading (30 wt% PANI). In these fibers,
PANI was finely dispersed around PPD-T domains,
and was oriented parallel to fiber axis. Fibers
containing 5 and 30 wt% PANI had an electrical
conductivity in the range of 1024 and 0.1 S/cm,
respectively.
Sometimes it is useful to decrease the concen-
tration of PANI in H2SO4. Thus, Ogurtsov and
Pokhodenko [231] used such solution to prepare a
PANI/Nylon 6 composite with a low percolation
threshold (0.03–0.07 wt% of PANI). They found that
the decrease of ionic strength of the solution when
lowering the PANI concentration led to an unwrap-
ping of the macromolecular chains. An increase of an
interface surface tension and reduction of the
percolation threshold accompanied this process.
It should be noted that it is more convenient the
use of liquid organic acids than sulfuric acid as
solvents for PANI solution blending, due to ease of
handling and solvent removal. For example, Abra-
ham et al. [232] prepared free standing, lustrous
and flexible blend films of PANI and Nylon 6 at
various weight ratios by casting from homogeneous
solutions in formic acid. The maximum conduc-
tivity of the films was about 0.2 S/cm correspond-
ing, for a weight ratio of 0.5 (w/w) PANI and
Nylon 6. A simultaneous TGA–DTA scan revealed
that the melting temperature of PANI/Nylon 6 was
slightly reduced, and an X-ray diffraction pattern
indicated that the crystal structure of Nylon 6 in
the blend was retained on modification. These
results demonstrated the absence of any substantial
interaction between the two components of the
blend [232], in contradiction with the numerous
data on interaction of Nylon 6 matrix polymer with
PANI discussed above.
Anand et al. [233] also used formic acid for the
preparation in a solution of blends of PANI
derivatives POT and PMT with 10–90 wt%
PMMA. The blend was precipitated by the addition
of the formic acid solution to water (non-solvent).
The thermal stability of the blends was higher than
that of individual POT-HCOOH and PMT-HCOOH
salts, and the conductivity of POT (30)–PMMA
(70) and POT (30)–PMMA (70) blends was close to
1026 S/cm.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1725
The strong acid MSA was used by Su et al. [234] as
the solvent to prepare blend films of PANI and poly(4-
vinylpyridine) (P4VP), with PANI loading from 100
to 50 wt%. It was suggested that MSA as blending
solvent formed hydrogen bonds with both PANI and
P4VP. Dry PANI-MSA/P4VP films prepared by
vacuum distillation had conductivities in the range
of 2.9 £ 1023–4.6 £ 1021 S/cm. The blend with
80 wt% of PANI showed an interesting elliptical
flake morphology, in contrast to the spherical particle
morphology observed for other blends.
Recently, Adams et al. [235,236] developed a new
acid-solution processing route for preparation of
highly conductive PANI films and fibers. It comprises
the use of AMPSA as both the protonating acid and
the solvating group, and dichloroacetic acid (DCAA)
as the solvent. The AMPSA content was varied so that
between 30 and 100% of the nitrogen sites on PANI
could be protonated. A modification of this route
involving the solution blending of PANI with a new
multifunctional dopant DEHEPSA and PMMA in
DCAA or difluorochloroacetic acid resulted in
flexible conducting composite films with a low
percolation threshold (much below 1 wt% of PANI)
[199,237]. The possibility to prepare such composites
was based on the fact that the diesters of 5- or 4-
sulfophthalic acids improve PANI solution processi-
bility [237–239]. Polyaniline protonated with these
acidic esters was soluble in chloroform, diethylk-
etone, hexafluoro-2-propanol, m-cresol and dichlor-
oacetic acid. Olinga et al. [237] found that the use of
DEHEPSA together with DCAA as solvent led to
PANI-DEHEPSA films with conductivity of 180 S/
cm. These films demonstrated metallic-like behavior
down to 220 K.
Usually, PANI doped with some sulfonic acids and
processed from a solution exhibits poor mechanical
properties. The introduction of sulfonic group in the
classical plasticizers benzenedicarboxylic acid die-
sters as the dopant resulted in PANI that exhibited
good mechanical properties, excellent flexibility and
much lower glass-transition temperature, as compared
to PANI doped with other protonating agents. Blends
of PANI-DEHEPSA/PMMA also had better mechan-
ical properties compared to PANI-CSA/PMMA cast
from m-cresol [238]. For example, the elongation at
break increased from a few percent (the CSA case) to
at least 40% (the DEHEPSA case). It should be noted
that upon casting, the DCAA solvent was efficiently
removed from the polymer matrix, so that the
resulting blends did not release solvent with age, as
tends to occur with blends cast from m-cresol.
3.1.5. Blends prepared of joint PANI base and
common polymer solutions in NMP
Angelopoulos et al. [147] discovered that the EB
form, i.e. PANI sample deprotonated by treatment in
alkaline solution, readily dissolves in NMP. A low
molecular weight fraction of the undoped form is also
partially soluble in DMF [240,241]. Tzou and
Gregory [242] found that EB also easily dissolves in
N,N0-dimethylpropylene urea (DMPU). Specifically,
its 20 wt% solution in DMPU is much more stable in
time to a gelation process than its much less
concentrated solutions in NMP. Nevertheless, NMP
is still the most frequently used solvent for the
treatment of the emeraldine base.
The solubility in NMP became the basis to form
different protonated PANI composites by mixing
corresponding EB solutions with solutions of dopant
and matrix polymer. For instance, Yin et al. [243]
prepared a conducting PANI/PC composites by
dissolution of the emeraldine base, PC (Lexan 141)
and TSA in separate NMP portions to give 4, 10 and
5 wt% solutions, respectively. These solutions were
mixed to form a uniform solution, followed by casting
on a glass plate and drying. The conductivity of the
composite films was maintained in a wide range from
10218 to 0.01 S/m and was anisotropic, with the
conductivity parallel to the film surface larger than
that perpendicular to the surface. The percolation
threshold of 0.26 wt% for the parallel conductivity
was much less than the 9.5 wt% threshold for the
perpendicular conductivity, with corresponding criti-
cal exponents of the percolation law of 2.0 and 3.0,
respectively. These features were explained by quite
different morphology of the composite film in the two
directions. The authors [243] found that conductivity
of this composite also depended on the temperature
treatment, and was stable up to 160 8C only at high
PANI-TSA concentration (lower than that of pure
PANI-TSA, see below and Ref. [80]).
The marginal thermostability of the conductivity of
composites prepared through NMP solutions can have
some causes may be related to residual NMP retained
in the blends and composites, even when careful
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531726
drying [244]. This can lead to a PANI reaction with
NMP at elevated temperatures, as demonstrated by
Afzali [245].
NMP is very convenient in the preparation of
conducting composites of PANI with polymeric acids.
For example, Fu et al. [246] used this solvent to
realize protonation of PANI base with lightly
sulfonated PS. They showed that a relatively low
concentration of sulfonic acid groups in the polymer
(5.3 mol%) was sufficient for doping PANI, and
promoting a solubility of the resulting macromolecu-
lar complexes. Hu et al. [247] reported electrically
conducting PANI–poly(acrylic acid) (PAA) blend
coatings. The samples showed moderate electrical
conductivity, about 1025 S/cm in the range of the EB:
PAA molar ratio from 0.25 to 1. Immersion in
aqueous HCl produced an increase in conductivity of
two to three orders of magnitude, and a slightly
improved thermal stability. The loss of conductivity in
the both cases at temperatures higher than 130 8C was
attributed to HCl evaporation and/or the decompo-
sition of carboxylate groups of PAA [247].
Moon and Park [248] have prepared conducting
composites of PANI with copolymeric acids such as
poly(methylmethacrylate-co-p-styrenesulfonic acid)
(PMMA-co-SSA), poly(styrene-co-p-styrenesulfonic
acid) (PS-co-SSA), and poly(methylmethacrylate-co-
2-acrylamido-2-methyl-1-propanesulfonic acid)
(PMMA-co-AMPSA). Emeraldine base, PMMA-co-
SSA and PS-co-SSA were dissolved in NMP, and the
PMMA-co-AMPSA was dissolved in DMF, selected
as a better solvent thanNMP. The conductivity of these
composites was investigated as a function of the acid
content in the copolymeric acid dopants. It was found
that even if the fixed mole ratio of acid to aniline was
kept at the excessive value of 1:1 in the PANI
composites, the conductivity of the copolymeric
acid-doped PANI decreased with decreasing of acid
content in the copolymeric acid chains. This was
attributed to the non-acidic units in the copolymeric
acids, preventing doping of PANI by adjacent acid
groups. The PANI/PMMA-co-SSA composites
showed the highest conductivity, up to 0.001 S/cm,
up to about two order of magnitude higher than that of
the PANI/PMMA-co-AMPSA composites. The lack
of conductivity of the PANI/PMMA-co-AMPSA
composites was explained by the inefficient doping
ability of the bulk AMPSA groups. On the other hand,
the higher conductivity of the PANI/PMMA-co-SSA
composites in comparison with PANI/PS-co-SSA was
explained by hydrogen bonds formed between the
carbonyl groups in PMMA and the imine groups in
PANI, which could hinder phase separation and induce
more homogeneous mixing and efficient doping.
Sixou et al. [249] presented a comprehensive study
of the transport properties of PANI(EB)/Nafion and
(lithiumtrifluoromethanesulfonimide PANI complex)/
PEO blend films cast from NMP solutions. They
considered electronic transport processes in the
(PANI complex)/PEO and PANI(EB)/Nafion blends
in relationship to the organization of the PANI phase
and the PANI protonation levels. Specifically, hop-
ping and tunneling processes and doping heterogene-
ities of PANI were taken into account, and the
transport processes were explained in the framework
of the hopping model between highly conducting
PANI clusters. Concerning (the PANI complex)/PEO
blends, the average doping level of PANI did not
depend on the composition. An increase of the PEO
concentration resulted in a decrease of the fraction of
the highly conducting regions in the PANI pathway.
In the PANI(EB)/Nafion blends, the situation was
quite different, due to the performance of Nafion as
dopant. While the volume content of PANI was
increased, it appeared that the average doping level of
PANI decreased, and the conductivity went through a
maximum and than decreased. It was shown that the
maximum resulted from the competition between two
opposite effects of composition on the blend conduc-
tivity: (i) an increase due to scaling law of classical
percolation theory, and (ii) a decrease coming from
decreasing intrinsic conductivity of the percolation
network that was induced by lower the doping level of
PANI. The conductivity of the PANI/PEO composite
reached values 0.004 and 0.08 S/cm at 0.15 and
0.5 vol% of PANI, respectively, and the correspond-
ing maximum conductivity of the PANI/Nafion
system was 0.1 S/cm at 0.2 vol% of PANI.
3.2. Thermally processible PANI blends and
composites
Thermally processible conducting polymer blends
and composites are more practical in industrial scale
than solution processed system. As a consequence,
this stimulates researchers and manufacturers to their
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1727
development. Three main approaches exist to produce
such materials. The first is realized through the
mechanical dispersing of infusible conducting poly-
mers in melt thermoplastic matrixes to achieve
conventionally moldable or extrudable conductive
composites. The second is the development of melt
processible electrically conductive polymers. The
third combines the preparation of a mixture of PANI
in a dispersion with a thermoplastic polymer solution
or dispersion (considered above in Section 2),
followed by the separation of the mixture and its
melt treatment (compression molding, extrusion, etc.)
3.2.1. Composites with infusible PANI
The principal requirements for use of a PANI as an
infusible component in a composite are easy dis-
persion in thermoplastic matrix polymers and suffi-
cient thermal stability in processing and operation
conditions. Particles of PANI produced by standard
techniques through oxidative aniline polymerization
in an inorganic acid water solution have a high surface
tension, resulting in their tendency to aggregate, and a
lowered specific surface. After drying, large aggre-
gates of PANI particles are formed, with sizes up to
several hundreds of microns. Replacing inorganic
acid dopants by functionalized protonic acids
improves the situation. Thus, Shacklette et al. [80]
have developed PANI compositions having a surface/
core dopant arrangement in which a dopant at or near
the surface of PANI particles (up to the depth of about
50 A in the skin) is different from for a dopant at or in
the core of these particles (about 50 A from the
surface). This structure is motivated by the use of
the dopant in the skin to impart high conductivity to
the PANI particle surface and in the core to promote
thermostability and processibility. In addition, such a
structure considerably decreases the surface tension of
the particles, which then form aggregates that are
easily demolished and dispersed in the melt of a
thermoplastic polymer.
Shacklette et al. [6] found that the commercial
form of PANI doped by TSA-Versicone consists of
aggregates with a basic morphological feature,
characterized as spheres within spheres. The average
powder grain, with a dimension of about 50 mm,
comprised a collection of small spheres (,1 mm in
diameter). In turn, the latter comprised smaller
spheres with sizes from ,0.05 to ,0.2 mm, built
from still smaller primary particles, ,10 nm in size.
Pelster et al. [250] concluded that the primary 10-nm
diameter particle has an 8-nm metallic core sur-
rounded by a,1.6 nm amorphous non-metallic shell.
Basing on small-angle X-ray scattering data later,
Wessling [251] suggested that one primary particle
(10–15 nm) might consist of ,20 individual mol-
ecules, folded to a diameter of 3.5 nm, to form a
coherent metallic core. Lennartz et al. [252] shown
that these PANI-TSA primary particles agglomerate
to around 50 nm aggregates in a PMMAmatrix. These
particles are considered to be the hyperstructure
which formed secondary particles of ,100 nm in
polymer matrices [251,252].
Versicone was dispersed in thermoplastic PVC,
PETG, and PCLU using conventional compounding
in a Brabendere mixer [6]. Specifically, intensive
mechanical mixing of Versicone and molten PVC
resulted in PANI particles 100–200 nm in size.
Percolation curves of these composite obeyed the
standard function s ¼ s0ðw2 wcÞt; derived from the
random percolation theory, where wc is the critical
volume fraction (content) of the conductive filler
necessary to achieve percolation and s0 is the intrinsic
conductivity of the filler. The value of the exponent t
is generally thought to be universal, with a theoreti-
cally predicted value near 2. The values of these
parameters for several composites (Table 2) demon-
strate their dependence on the matrix [6]. As one can
see from this table, the value of wc was significantly
lower in the Versicon composites than would be
expected for a random dispersion of the particles. This
suggests that the dispersed PANI particles partially
reaggregated at some point during processing and
molding, due to PANI incompatibility with the matrix
polymer. This incompatibility indicates a mismatch in
surface energy or solubility parameter causing
Table 2
Percolation above the critical concentration (of the Versicon
composites)
wc (vol%) s0a (S/cm) t
Random filling (3D) 0.15–0.30 – 1.6–2
PANI in PETG 0.062 93 2.8
PANI in PCL 0.046 292 1.9
a s0 , 6 S/cm for Versicon [80], 1–20 S/cm for unoriented bulk
PANI and 100–300 S/cm for pure oriented samples of PANI [6].
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531728
the PANI particles to be driven from regions of the
melt and collected at the periphery of matrix domains,
forming one and two-dimension aggregated structures
[6]. Such a distribution of conductive particles leads
to dramatically lower critical volume fractions in
systems for which the average dimension of insulating
domains is much larger than that of the conductive
particles [253]. Therefore, the excellent percolation
results obtained for these composites are derived from
the small primary PANI particle size.
Unexpectedly, as may be seen in Table 2 (and its
footnote) the calculated s0 was significantly higher
for the blends than for Versicon or unoriented bulk
PANI [6]. Similar behavior is observed for other
mixtures. Thus, when investigating the electronic
transport properties of PANI-TSA/PMMA and PANI-
TSA/PVC blends Kaiser et al. [254] found that
blending PANI with PMMA and PVC increased the
conductivity, especially at lower temperatures (Fig. 4).
This increase was ascribed to lessening of insulating
barriers around PANI particles in these blends. The
temperature dependence of the conductivity in PANI
blends was well described by a series combination of
quasi-1D metallic resistivity and tunneling (between
small metallic islands). However, it should be stressed
again that unlike these PMMA and PVC cases and
the results in Table 2 [6], blending PANI with
heterochain copolyester PETG (analog of PET, with
the ability to form hydrogen bonds with PANI [115])
gave reduced conductivity as expected from general
considerations [254,255]. The striking contrast
between the conductivity for PANI-TSA/PETG
blends and PANI-TSA/PMMA composites is consist-
ent with a picture of tunneling between metallic
particles separated by non-metallic barriers. The
conductivity of the PANI-TSA/PMMA blend with
60 wt% PMMA exceeded that of pure PANI-TSA at
all temperatures. By contrast, the conductivity of a
PANI-TSA/PETG blend with 60 wt% of the non-
conducting polymer (PETG) was several times less
than that of the unblended PANI-TSA. Near room
temperature, unblended PANI and PANI-TSA/
PMMA blends both showed a change to metallic-
like temperature dependence of the conductivity,
whereas this did not occur for the PANI-TSA/PETG
blends. The approximate linearity of the logarithm of
conductivity as a function of 1=T1=2 showed that the
conductivity over a wide temperature range was
generally consistent [254,255] with the usual form
[256] for granular metals
s ¼ s0 expð2ðT0=TÞ1=2Þ ð3Þwhere s0 and T0 were constants.
However, the conductivity above 250 K deviated
strongly from Eq. (3) for the more highly conducting
blends, and changed to a metallic sign for the
temperature dependence in PANI-TSA and PANI-
TSA/PMMA blends. For these cases the composite
expression for conductivity was found to be [254]
s21 ¼ r expð2Tm=TÞ þ t expððT0=TÞ1=2Þ ð4Þwhere the coefficients r and t determined the
magnitude of the metallic and tunneling resistivity
terms, respectively, but depended on morphology in a
complex fashion. Instead of the coefficient r and t; the
fits were made in terms of the conductivity sð300Þ at300 K and fraction m of the resistance at 300 K
arising from the first (metallic) term in Eq. (4), given
by Ref. [255]:
m ¼ rsð300Þexpð2Tm=300Þ ð5ÞThe value of sð300Þ determined the overall
conductivity magnitude, while m indicated the extent
of the reduction in slope near room temperature
Fig. 4. Themperature dependence of conductivity of two PANI-
TSA/PMMA blends (with 60 and 67% PMMA) and unblended
PANI-TSA [255]. Reproduced from Subramaniam et al. by
permission of Solid State Commun 1996;97:235. q 1996 Elsevier
Science Ltd, Oxford, UK.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1729
(whether or not a change to metallic sign occurred
also depended on the value of T0). The resulting fit
parameters are listed in Table 3.
The value of Tm represents the energy of phonons
with wave vector spanning the Fermi surface of the
highly anisotropic metal. Since Tm is not determined
accurately by the data, a value of 2000 K was taken
for all samples. The fitted values of m in Table 3 show
the largest metallic contribution for PANI-TSA and
PANI-rich blends, and small values for the low
conductivity PANI-TSA/PETG blends. The values of
T0 in the tunneling term are much smaller for the
PMMA blends, reflecting the much smaller decrease
of conductivity in these samples as the temperature
decreases. The values of T0 for the PETG blends show
only a relatively small change from that for unblended
PANI, suggesting that the PANI particles retained
their original properties to a greater extent than in the
PMMA blends. The thermopower of PANI and of all
these blends was small, and (apart from PANI at low
temperatures) increased with temperature. This
remarkable behavior of the thermopower of all blends
resembled metallic diffusion thermopower, in contrast
with the huge difference in conductivity [255].
Srinivasan et al. [257] presented additional
evidence, based on a Dysonian line shape in ESR
studies, for the metallic nature of PANI-TSA
(Ormeconw) and its (33 wt%) blend with PMMA,
prepared by a dispersion technique under shear
condition in melt phase. They showed that at low
temperatures the line shape became symmetric and
Lorentzian when the sample dimensions were small in
comparison with the skin depth. It was also found
that the unblended PANI had a much stronger
temperature dependence of the conductivity than
the PANI(33 wt%)–PMMA(67 wt%) blend. For this
blend, the activation energy of the dependence of
conductivity on temperature increased as decreasing
T below 1 K, showing behavior for the metallic side
of the metal–insulator transition. By contrast, the
activation energy decreased with decreasing T in the
same temperature range for unblended PANI, show-
ing behavior for the insulating side of the
transition. The authors claimed that this showed that
the metal– insulator transition appeared only for
materials prepared with a dispersion step in the
processing [257].
The significance of the dispersion of the PANI
phase for conductivity properties of its blends and
their dependence on a matrix polymer was demon-
strated by Zilberman et al. [258,259]. They investi-
gated Versicone melt-mixed blends with
thermoplastic polymers such as PS, PS plasticized
with DOP, PCL, CoPA, LLDPE and LDPE. The
blending temperature was chosen depending on the
matrix polymer. Thus, blend temperatures were given
by PS or LLDPE at 180 8C, plasticized PS at 150 8C,CoPA at 165 8C, LDPE at 130 8C and PCL at 70 8C.The results showed that the blend morphology and the
level of interaction between components of the blends
strongly affected the electrical conductivity of the
blend, as may be seen from the dependence of the
electrical conductivity as a function of PANI-TSA
content given in Fig. 5. These data showed that
percolation began in the range of,5–10 wt% PANI-
TSA for heterochain polar polymers (PCL, CoPA)
and plasticized PS-DOP blends. By contrast, PANI-
TSA blends with non-polar carbochain polymers
(LLDPE, LDPE, PS) were conductive only at PANI-
TSA loadings higher than ,30 wt%. SEM and TEM
studies of the blend morphology displayed large
agglomerates (5–50 mm) of the PANI particles within
LLDPE and PS, indicating the absence of a continu-
ous network of PANI even at 20 wt% PANI-TSA. In
contrast, the blends of CoPA or PS/DOP exhibited
dispersed small PANI particles (0.1–0.5 mm). More-
over, the PANI-TSA particles in the plasticized PS/
DOP matrix were smaller than those in the CoPA
matrix and even more, in the PCL-based blend with
the lowest percolation threshold (,5 wt% PANI-
TSA) only a few very small particles were registered.
As a consequence, the higher conductivity at rela-
tively low PANI-TSA content in PCL and CoPA was
Table 3
Parameter values for the fits of Eq. (5) to the conductivity data for
the PANI blends [254]
Blend s (300, S/cm) T0 (K) m
PANI-TSA(40%)/PMMA 30 130 0.09
PANI-TSA(33%)/PMMA 13 60 0.11
PANI-TSA(40%)/PETG 3.6 770 0.11
PANI-TSA(30%)/PETG 0.91 650 0.06
PANI-TSA(20%)/PETG 0.10 950 0.03
PANI-TSA(15%)/PETG 0.013 1350 0.02
PANI-TSA 18 1040 0.28
The value of Tm was taken as 2000 K.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531730
assigned to the higher levels of dispersability and
structuring of the PANI-TSA particles within the
matrix polymer. Thus, the PCL matrix in the PANI-
TSA(20 wt%)/PCL and PANI-TSA(5 wt%)/PCL
blends showed spherulitic crystallization, in which
the spherulites in the 5/95 blend were similar to those
of the neat PCL, and larger than those obtained for the
20/80 blend. The included PANI-TSA particles were
located around the spherulites in the amorphous
regions [258,259], as usual for additives in semi-
crystalline polymers. Hence, the doped PANI network
located within these regions, leading to reduction of
the percolation concentration. This accounts the
significantly lower percolation threshold for the
PCL-based blends (5 wt%) in comparison with that
for the amorphous-matrix-based blends, PANI-TSA/
CoPA (10–15 wt%) and PANI-TSA/PS-DOP
(10 wt%) [258,259]. These interpretations accord
with those of Shacklette et al. [6] discussed above
on the effects of mismatches in surface energy and in
solubility parameter on the distribution and reaggre-
gation of PANI particles in a matrix polymer.
Indeed, interaction of the matrix polymer with the
conducting polymer effects dispersion of the conduct-
ing phase in the matrix during MP, and higher PANI
fracturing is observed for matrices interacting more
strongly with PANI, (at similar matrix viscosity and
shear level), due to better interphase shear stress
transfer [258,260]. Zilberman et al. [258,259] showed
that this occurred in systems with components
with similar solubility parameters. Specifically,
the CoPA-based blends were compatibilized better
than blends containing LLDPE and PS. In the first, the
solubility parameters of the components were similar,
whereas the solubility parameters of the matrix
polymers were well below than that of PANI-TSA
for LLDPE- and PS-based blends (Table 4). As a
result the addition of 20 wt% PANI-TSA to CoPA
increased its Tg by 4 8C, suggesting specific inter-
actions in PANI-TSA/CoPA system, unlike the
addition of 20 wt% PANI-TSA to LLDPE or PS,
which had negligible affect on Tg: Such specific
interaction may include hydrogen bonds between the
hydrogen of the amine nitrogen of PANI and
Fig. 5. Electrical conductivity versus PANI-TSA content of polymer/PANI binary blends. Compiled from Ref. [258]. Zilberman M,
Siegmann A, Narkis M. J Macromol Sci, Phys 1998;B37(3):301 and Ref. [259] Zilberman M, Siegmann A, Narkis M. J Macromol Sci
Phys 2000;339(3):333, by courtesy of Marcel Dekker, Inc.
Table 4
Solubility parameters of PANI and polymer matrix calculated
without taking a specific interaction into account [257]
Polymer Solubility parameter (J/cm3)0.5
PANI-TSA 23.9
PCL 17.8
CoPA 24.2
LLDPE 16.8
PS 19.5
DOP 20.9
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1731
the carbonyl oxygen in polyamides. Compatibility of
PANI-TSA with insulating PS was improved by the
addition of plasticizer (15 wt% of DOP) to PS,
resulting in conductive PANI-TSA/PS blends. This
suggested that the DOP addition increased the
solubility parameter of PS towards that of PANI-
TSA. On the other hand, an additional factor could be
migration of DOP into the PANI phase during
blending, affecting the PANI rheological behavior.
Furthermore, the addition of more PANI-TSA
(20 wt%) quantity to PS-DOP increased Tg of PS by
5 8C. The effect of DOP content on the electrical
conductivity of PANI-TSA(20 wt%)/PS-DOP may be
seen in Fig. 6. Zilberman et al. [259] believed that
DOP acted not only as plasticizer, but also as a
compatibilizer, to improve the PANI–PS interaction,
giving PANI-TSA/PS-DOP blends that conduct at
relatively low PANI-TSA content.
Zilberman et al. [259] investigated the conductivity
and morphology of PANI-TSA/CoPA/LLDPE and
PANI-TSA/PS-DOP/LLDPE ternary blends. They
found the important fact that the doped PANI
preferentially located in one of the phases, due to
increased interactions between PANI and the preferred
polymer. Thus, in the case of the PANI-TSA/CoPA/
LLDPE blends, the PANI phase preferentially located
in the CoPA to give an effective PANI content in the
CoPA phase higher than its nominal content in the
blend. The system specificity led to a double-percola-
tion phenomenon in the ternary blends containing
10 wt% PANI (Fig. 7) resulting in high conductivity
for the blend based on CoPA/LLDPE 30/70.
As one may see from Fig. 7, binary blends based on
CoPA and LLDPE at 10 wt% PANI were insulating. It
might be expected that if most of the PANI particles
were located at the CoPA/LLDPE interface, a very low
percolation threshold would be observed. Energy
dispersive spectroscopy sulfur mapping of the fracture
surfaces of the blends showed that about 90%of PANI-
TSA was located within the CoPA phase, with
remainder within the LLDPE phase or the CoPA/
LLDPE interphase. Therefore, the authors [259]
concluded that conductivity of the PANI-TSA/CoPA/
LLDPE blend was determined mainly by the PANI
content within the preferred phase, its mode of
dispersion, and the conducting network structure
created. The solubility parameter of PANI-TSA (see
Table4 andRef. [258])was found tobe similar to that of
CoPA, and to be much higher than that of LLDPE, so
that only a portion of thePANIparticles remained at the
CoPA/LLDPE interface. As with the PANI-TSA/
CoPA/LLDPE blends, PANI-TSA/PS-DOP/LLDPE
blends also consisted of polymers which were compa-
tible (PS-DOP)and incompatible (LLDPE)withPANI-
TSA.Hence, onemight expect a similar behavior of the
two systems. However, a different behavior was found
for their electrical conductivity: the conductivity level
of blends containing 20 wt% PANI slightly decreased
with increasing LLDPE content, whereas the blends
containing 10 wt% PANI were all insulating, like
Fig. 6. The electrical conductivity (A) and Tg (W) as a function of DOP content for the PANI-TSA(20 wt%)/PS-DOP blends [259]. Reprinted
from Zilberman M, Siegmann A, Narkis M. J Macromol Sci Phys 2000;B39(3):333 by courtesy of Marcel Dekker, Inc.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531732
the corresponding binary blends. It was shown that
PANI was located mainly within the PS-DOP phase,
with only a small quantity in the LLDPE.This behavior
was expected from the calculated solubility parameters
shown in Table 4, i.e. PANI tended to locate within the
more compatible matrix polymer. In this case, the
highly effective PANI content in the (PS-DOP) phase
did not generate electrical conductivity for PANI-TSA/
PS-DOP/LLDPE ternary blends containing 10 wt%
PANI-TSA. This phenomenon was explained by some
migration of DOP from PS, which left PS less
plasticized, and less compatible with PANI [259].
A study of polymer composites with PANI-TSA
produced without use of special expedients weaken-
ing the interaction among the ‘primary’ particles of
PANI, used PANI-HCl prepared by a standard
technique [261,262]. The product was neutralized
with NH4OH and then redoped with TSA to be used in
melting blends with thermoplastic polymers. Specifi-
cally, Mitzakoff and De Paoli [261] prepared blends
of PANI-TSA and engineering plastics (PET and
Norylw) by mechanical mixing at 260–270 8C and
5 min in a Haacke torque rheometer. The Noryl
employed in that work was a 1:1 blend of poly-
phenylene oxide and high impact PS. However, the
PANI-TSA particles used were too large (between
62–149 and 44–62 mm) to allow good percolation.
Nevertheless, despite this and the severe mixing
conditions, the conductivity for blends with 5% of
PANI-TSA, stabilized at ca. 1025 and 8 £ 1027 S/cm
for the PANI-TSA/PET and PANI-TSA/Noryl blends,
respectively. The lower conductivity for using Noryl
compared with PET was explained by differences in
the resistivity of these polymers, 10218 against
10216 S/cm, respectively. Based on an investigation
of mechanical properties of the blends, Mitzakoff and
De Paoli [261] made the important conclusion that the
acidic dopant of PANI caused hydrolysis of the ester
bonds of PET, producing a hard and brittle material
that hindered its application. However, the PANI-
TSA/Noryl blends had good mechanical properties,
with conductivity in a range useful for the production
of plastic parts able to dissipate electrostatic elec-
tricity. The results showed that mechanical properties
improved both with decreased PANI loading and
better homogenization. The values of the Young’s
modulus ðEÞ for the blend changed with PANI loadingin the range of 1.1–1.5 GPa. Eq. (6) was proposed to
correlate the two variables analyzed with the
elongation at break ð1bÞ1b ¼ 9:99þ 0:097R2 1:78P ð6Þwith 1b given in percentage; R; the rotor speed in rpm;
P is the PANI concentration in percentage. From
Fig. 7. Electrical conductivity versus LLDPE content for PANI-TSA/CoPA/LLDPE ternary blends containing 10 and 20 wt% PANI-TSA [259].
Reprinted from Zilberman M, Siegmann A, Narkis M. J Macromol Sci. Phys 2000;B39(3):333 by courtesy of Marcel Dekker, Inc.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1733
Eq. (6) it is possible to estimate the elongation at
break of a PANI-TSA/Noryl blend containing
between 1 and 5 wt% of PANI-TSA, processed for
3 min at 260 8C.
Faez et al. [262] described the preparation and
the electrical, mechanical, thermal and morphological
characteristics of a conductive blend of the elastomer
EPDM and PANI-TSA. Polymer mixtures were
prepared in a similar manner to that described above
for PANI-TSA/PET and PANI-TSA/Noryl blends, but
at 150 8C and with different a mixing time and PANI-
TSA concentration. Specifically, the mixing time was
12 min for 0.5–10 phr (parts per hundred) of PANI,
16 min for 20–30 phr and 20 min for 40–50 phr.
Particle sizes were between 100 and 200 mesh (about
75–150 mm). The samples were vulcanized at 175 8C
and 5 MPa pressure for 15 min using dicumylper-
oxide. The results of the prepared blend testing are
presented in Table 5. As one can see, there is an initial
increase of the elongation at break for 5 phr PANI-
TSA content, followed by a decrease at higher PANI-
TSA content. The modulus showed a slow increase
between 5 and 30 phr PANI-TSA, and an abrupt
increase for the mixtures with 50 phr PANI-TSA. This
was attributed to the rigidity of PANI acting as a
reinforcing filler and changing the viscoelastic
behavior of the rubber to that of a rigid material
[262]. PANI-TSA contributed also to an increase
in the rubber cross-linking density as determined in
the swelling measurements (gel fraction, GF increase
in Table 5). At the same time, there was no variation
in Tg of the EPDM phases, suggesting that the
mixtures were immiscible. By changing the content of
PANI-TSA and controlling the mixing parameters it
was possible to produce vulcanized conductive
materials with elastomeric properties. The composite
conductivity increased continuously with PANI-TSA
content and at 50 phr seemed to reach a plateau at the
level of ,1026 S/cm [262].
3.2.2. Polymer blends and composites with fusible
PANI
The discovery of counter-ion induced solubility
[2,263] appeared to be the base of resolution of the
problem of imparting the MP capability to PANI
through the use of some functionalized sulfonic
acids, e.g. DBSA or phosphoric acid aliphatic and
aromatic diesters as doping agents [7,221–223,
264–266]. The conventional method for doping
EB is mixing with a functionalized protonic acid in
an appropriate solvent. However, a doping process
without solvent use by mechanical mixing EB with
DBSA [7,267] or phosphoric acid diesters [223] etc.
is more practical. Using this approach, Ikkala et al.
[7] developed conducting polymer blends by
conventional melt mixing of thermoplastic bulk
polymers with Neste Complex, a proprietary con-
ducting polyaniline composition PANI(DBSA)y. The
percolation threshold for conductivity was observed
at a low weight fraction of the PANI(DBSA)y,
differing with a matrix. In particular, they showed
that the acceptable for practice level of the electrical
conductivity of blends of Neste Complex could be
obtained with the polymer matrixes of different
origin:Table 5
Mechanical properties and gel fraction of pure rubber and blends as
a function of polyaniline concentration in phr: Young modulus ðEÞ;strain at break ðsbÞ; elongation at break ð1bÞ and gel fraction (GF)
[261]
PANI concentration
(phr)
E
(MPa)
sb
(MPa)
1b £ 1022
(%)
GF
0 3 5.4 6.0 0.94
5 3 7.6 9.4 0.96
10 5 6.6 6.4 0.96
20 9 6.0 3.5 0.96
30 14 6.5 2.8 0.97
40 13 8.0 2.0 0.99
50 26 7.9 1.0 0.99
Material PANI(DBSA)y/log s
(wt%)
High density polyethylene 3.2/24.33; 7.8/22.12
LDPE 2.4/24.31; 4.3/21.67
Polypropylene 1.9/24.33; 7.6/21.77
PS 2.8/26.23; 7.9/22.05
Impact modified PS 4.0/21.78; 7.2/21.05
PVC 1.5/24.37; 2.4/21.22
Poly(styrene–ethylene/
butylene–styrene)
1.1/27.25; 1.8/20.13
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531734
These composites were proposed for applications
such as electrostatic dissipation (ESD), static
discharge and EMI shielding, which require
conductivities of approximately 1025–1029 S/cm
for ESD and .1 S/cm for EMI. Ikkala et al. [7]
concluded that the ESD level conductivity could be
achieved by application of only a few percents of
Neste Complex. Even conductivity levels near those
required in EMI shielding have been achieved in
some cases.
Ahlskog et al. [268] found the doping reaction of
PANI with mechanically mixed DBSA is a time-
dependent process, accelerated by heating. An extra
amount of DBSA yielded a plasticized melt
processible complex [269,270]. Specifically, a fully
doped state is obtained for a PANI:DBSA molar ratio
of 1:0.5. Use of an excess amount of DBSA led to
a decrease of Tg and a plasticizing effect, resulting in
easier MP (Fig. 8). Thus, Tg was,135 8C for a molar
ratio of 1:0.7, compared with the much higher Tg(,230 8C) observed without a DBSA excess (the
molar ratio of 1:0.5). It was found that in this molar
ratio the critical DBSA mole fraction to achieve the
essential plasticization was 0.7 [269,270].
According to Titelman et al. [271] the thermal
doping process includes the following main stages:
heating the blend, exothermic PANI-DBSA doping
reaction accompanied by a paste-to-solid-like tran-
sition, and plasticization of the resulting PANI/DBSA
complex by excess of DBSA. They showed that the
blends prior to thermal processing already consisted
of partially doped PANI particles, with a core/shell
structure. The core consisted of PANI (base)
and the shell of the PANI(DBSA)0.32 complex.
When the doping reaction was completed at
the paste-to-solid-like transition, further mixing did
not affect the complex composition, but led to
a reduction in conductivity. Levon et al. [267] used
X-ray studies to reveal a layered structure with 2.7 nm
spacing between layers for the complex in the
presence of excess DBSA, leading to enhanced
processibility. In the absence of excess DBSA, there
was no layered structure and the PANI(DBSA)0.5complex was therefore not heat-processible [271].
Zilberman et al. [258,272] also investigated a
conductive PANI-DBSA complex prepared by
a thermal doping process at the weight ratio
EB:DBSA ¼ 1:3 in Brabender plastograph at
140 8C for 5 min. It was used for melt mixing with
thermoplastic polymers (PS, PS plasticized by DOP,
LLDPE, CoPA) at temperatures that varied with the
matrix polymer (see above for the PANI-TSA case).
Naturally, the electrical conductivity of the blends
depended markedly on the matrix polymer (Fig. 9),
and on the compatibility of the components. This
resulted in the fact that unlike the PANI-TSA case
the blends of PANI-DBSA with heterochain CoPA
and carbochain LLDPE polymers were poorly
conductive even at 20 wt% PANI-DBSA, whereas
for case of aromatic polymer PANI-DBSA/PS
blends there was suitable conductivity, with a
percolation threshold at the 5 wt% PANI-DBSA;
the conductivity attained 5 £ 1024 S/cm for
PANI-DBSA(30 wt%)/PS. However, the actual
PANI-DBSA content was smaller than the nominal
Fig. 8. Themperature dependence of the storage modulus of PANI-
DBSA mixtures measured using (a) three-point bending and (b)
parallel plate geometry. The parameters show the mole fraction of
DBSA [270]. Reproduced from Vikki et al. by permission of Synth
Met 1995;69(1–3):253.q 1995 Elsevier Science Ltd, Oxford, UK.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1735
value (shown in Fig. 9) due to the presence of excess
DBSA. SEM micrographs of cryogenically fractured
surfaces of PANI-DBSA/PS, PANI-DBSA/CoPA,
and PANI-DBSA/LLDPE blends show large
domains of PANI-DBSA dispersed in the CoPA
matrix. This indicates that continuous networks of
PANI-DBSA could not be formed, even in a blend
containing 20 wt% PANI-DBSA, which, in conse-
quence, was practically insulating (,10210 S/cm).
Smaller particles of PANI-DBSA (0.5–2 mm) were
observed in a LLDPE-based blend with a higher
conductivity (,1028 S/cm). This difference was
explained by the compatibilizing effect of the
aromatic ring and dodecyl alkyl chain of DBSA,
promoting a better conducting network and smaller
particles in the hot-melt blending with non-polar
aliphatic PE rather than with polar CoPA. This
agreed with the results for the PANI-
DBSA(20 wt%)/PS blends for which very small
(0.1–0.2 mm) PANI-DBSA particles were observed.
The authors [272] supposed that this behavior was a
result of the high fracturing level of the PANI-
DBSA particles due to their high interaction with the
PS matrix. The calculated solubility parameter of
PANI-DBSA (20.8 (J/cm3)0.5) and its interaction
with the various matrix polymers (Table 4)
supported the electrical conductivity results (Fig. 9).
Perhaps because their components exhibited quite
similar solubility parameters, the PANI-DBSA/PS
blends appeared the most suitable systems among
those considered to obtain an electrical conductivity
high enough for use, while the PANI-DBSA/CoPA
blends are the least suitable ones. The DOP
plasticizer increased the solubility parameter of PS
towards that of PANI-DBSA, resulting in enhanced
dissolution of PANI-DBSA in the PS matrix during
MP, and in a slightly higher conductivity [272].
Faez and De Paoli [273] also used fusible PANI-
DBSA in a blend with EPDM (compare with the case
of infusible PANI-TSA, Section 3.2.1 and Ref. [262]).
Previously, to prepare this PANI-DBSA complex they
doped EB by three methodologies: (s) stirring EB for
240 h in a 1.5 mol/l solution of DBSA; (m) grinding in
mortar EB and excess DBSA in the 1:2 ratio and (r)
doping EB with excess DBSA in 1:2 ratio by reactive
processing in an internal mixer at 150 8C for 10 min.
Conductivity values were 1023, 1 and 5 S/cm for
PANI-DBSA doped by grinding in mortar, solution
and reactive processing, respectively. For PANI-
DBSA(s)/EPDM blends prepared by processing, all
EPDM was dissolved in cyclohexane. In this case the
PANI-DBSA complex did not contain excess DBSA.
That resulted in a blend consisting basically of
particles dispersed in the matrix without cross-
linking. However, PANI-DBSA(m)/EPDM and
PANI-DBSA(r)/EPDM blends showed partial
Fig. 9. Electrical conductivity versus PANI-DBSA content for a PANI-DBSA/PS blend series and for various matrix polymer/PANI-DBSA
(80/20) blends [272]. Reproduced with permission from Zilberman et al. J Appl Polym Sci 1997;66(2):243 q 1997 JohnWiley & Sons limited.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531736
insolubility of the EPDM phase. This behavior
indicated some kind of cross-linking and physical
entanglement or chemical reaction. Conductivities of
the PANI-DBSA/EPDM blends were much less than
that of the PANI-DBSA complex, but increased
linearly with PANI-DBSA content until ,30 wt%,
reaching 8 £ 1027, 2 £ 1026 and 5 £ 1026 S/cm for
DBSA(s)/EPDM, PANI-DBSA(m)/EPDM and PANI-
DBSA(r)/EPDM blends, respectively. Later, these
authors [274] found that the use of similar DBSA and
EPDM concentrations gave PANI-DBSA(r)/EPDM
blends with the higher conductivity (from 1023 to
1021 S/cm). Faez et al. [275] showed by SEM that the
morphology of PANI-(DBSA)3(r)/EPDM blends
undergoes significant changes during mixing. For
example, initially very compact and hard agglomer-
ates of PANI-DBSA decrease in size and acquire a
sponge structure with increasing mixing time. Faez
et al. [276] demonstrated the possibility to prepare
conductive PANI-DBSA/EPDM blends, formed
under similar conditions, but cross-linked in two
ways: by chemical method (using phenolic resin) or
electron-beam irradiation. The blends had different
mechanical and conductivity properties, dependent on
the cross-linking method.
A strong dependence of conducting PANI blend
properties on the composition and processing con-
ditions has also been demonstrated for melt mixed
PANI-DBSA complex with SBS rubber [277,278].
Leyva et al. [278] showed that the conductivity is
enhanced for blending at a higher temperature
(130 8C) in Haake internal mixer compared to the
blend compression-molded at 100 8C. However, a
highly cross-linked material was obtained at the
higher temperature. It should be emphasized that the
mechanical performance of the PANI-DBSA/SBS
blends was not good in comparison with pure SBS.
Thus, the ultimate tensile strength and elongation at
break of compression-molded at 100 8C samples
decreased from 21.0 MPa and 5200% to 9.5 MPa
and 3900% for pure SBS and its blend with 17 wt%
PANI-DBSA loading, respectively. Interesting XPS
N1s core-level spectra of the blends prepared in
different conditions demonstrated that MP of PANI-
DBSA in the SBS matrix promoted an additional
protonation level of the PANI chains.
Koul et al. [279] reported blends of conventional
thermoplastic ABS copolymer with PANI doped with
a specific ratio of mixed dopants, consisting of DBSA
and TSA at dopant ratios from 1:1 to 9:1. Blending of
this PANI with ABS was carried out in a Nuchen
Extruder at temperatures ranging from 180 to 190 8C.It was found the blends had the best conductivity
when PANI was doped by mixtures of DBSA and
TSA in 1:1 and 9:1 ratio [279]. Specifically,
conductivities for PANI-DBSA–TSA(1:1)/ABS com-
posites were 7.6 £ 1028, 8 £ 1027, 1.3 £ 1025 and
0.1 S/cm for 20, 30, 40 and 50 wt% of PANI-DBSA–
TSA, respectively. The lowest loading of PANI doped
with hybrid dopants in the molded conducting
composites might be effectively used for the dissipa-
tion of electrostatic charge. With higher loading a
shielding effectiveness of 60 dB at 101 GHz was
achieved, which suggested the conducting composites
as potential EMI shielding materials [279].
Paul and Pillai [280] have studied the synthesis of
new doping agents prepared from inexpensive natural
materials. They reported sulfonic acid derivatives of
3-pentadecylphenol derived from cardanol were
excellent plasticizing dopants, imparting thermal/
solution processibility to PANI. Specifically, there
were synthesized SPDP, SPDA and SPDPAA, used to
prepare freestanding hot pressed flexible films of
heavily plasticized, protonated PANI. Maximum
conductivity values 65 and 42 S/cm were obtained
for the PANI-SPDPAA and the PANI-SPDA films,
respectively, pressed at 140 8C. This is even better
than the conductivity values of 1–20 S/cm reported
for the melt processible PANI-DBSA [7,267,268].
The conductivity was comparatively less (10 S/cm)
for the PANI-SPDP film. It was shown that all these
protonated polymers were thermally stable up to
200 8C and, correspondingly were suitable for the
preparation of highly conducting blend films by MP.
The conductivity values obtained for PANI-
SPDA/PVC blends were higher than those for the
PANI-SPDPAA/PVC system (Fig. 10). This agreed
with SEM data showing a continuous conducting
network in the PANI-SPDA/PVC system, but PANI-
SPDPAA agglomerates evenly distributed in the PVC
matrix for PANI-SPDPAA/PVC blends. The authors
believe that PVC, being a highly polar polymer,
enhanced the blending and compatabilization with the
less polar SPDA-doped PANI, but not with the polar
SPDPAA-doped PANI. This suggestion seems to be
incorrect since a polar substance is more compatible
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1737
with other polar material than with a non-polar one.
An explanation of the observed difference may be
based on data on the polarity of the substrates, the
detailed structure of the blends and the solubility
parameters of the blend components. For example, the
last was successfully used by Zilberman et al. [258,
272]. Paul and Pillai [280] found that the tensile
strength of the blend PANI-SPDA/PVC decreasing
rapidly with increasing plasticized PANI content, so
that the blend containing 25 wt% of PANI had a
tensile strength of 6 MPa. The increase of Tgwith increasing the PANI content in the blend was
taken to indicate miscibility of polymer blend
components.
Alkyl and aryl phosphoric acid diesters also
constitute an excellent group of PANI dopants,
which not only render this polymer conductive and
solution processible, but also plasticized it to be
melt processible [221–223,264]. Thus, plasticized
PANI exhibited rheological parameters characteristic
of a Bingham liquid, with the viscosity decreasing
with an increase of the diester content [222].
Protonation of EB with DiOHP resulted in a heavily
plasticized mixture which could be thermally
processed to give free standing films, with conduc-
tivity exceeding 10 S/cm [4]. Polyaniline freestand-
ing films with enhanced conductivity (65 S/cm, as
for PANI-SPDPAA [280]) were prepared by hot
pressing of PANI protonated with DPHP in
chlorobenzene [264]. It was shown that blends
with excellent mechanical properties could be
prepared by hot pressing (160 8C) PANI-DiOHP/
PVC plasticized by DOP or PANI-DPHP/PVC
plasticized by tricrezylphosphate (TCP). The con-
ductive blends demonstrated a low percolation
threshold (e.g. 6 wt% for PANI(DPHP)0.5/PVC-
TCP) [223].
Ikkala et al. [281] showed that the addition of
compatibilizers such as selected esters of gallic acid
favored the formation of a continuous PANI network
in thermally processed PANI/polyolefin blends.
Specifically, for PANI(DBSA)0.5/polypropylene
a percolation threshold lower than 1 wt% of PANI
was observed. Later, Yang et al. [282] used esters of
gallic acid to prepare composites of PANI doped by
diesters of phosphoric acid. To this end, EB, the
dopant BEHP, LG and LDPE, mixed by grinding in a
mortar for about 10 min, were fed into an extruder
and processed for 10 min at 130 or 150 8C. This
resulted in composites with percolation thresholds
below 3 wt% for the both temperature. The proces-
sing temperature had a small influence on the
conductivity despite the fact that BEHP is not very
stable in PANI at 150 8C. The authors [282]
Fig. 10. Conductivity versus PANI content in the blends of (a) PANI(SPDA)0.5/PVC, (b) PANI(SPDPAA)0.5/PVC. Pressing temperature,
160 8C; pressing time, 15 min [280]. Reproduced from Paul and Pillai by permission of SynthMet 2000;114(1):27.q2000 Elsevier Science Ltd,
Oxford, UK.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531738
supposed that LG acted as a compatibilizer, which
significantly modified interactions between immisci-
ble LDPE and PANI. They assumed the following
mechanism of the composite solidification. During
processing the molten gallate dissolves protonated
PANI, facilitated by the long flexible alkyl chain in
both the PANI dopant and the compatibilizer.
Similarly, the alkyl substituents of the compatibilizer
facilitate its miscibility with LDPE. Upon solidifica-
tion of the composite the compatibilizer forms a
continuous network within the LDPE matrix. Within
this network, in turn, microphase separation occurs
between the conductive PANI and LG. This micro-
phase separation is governed by a strong interaction
(probably via hydrogen bonding) between the polar
part of the compatibilizer molecule and PANI. Yang
et al. [282] supposed an existence of a double-
percolation network of the compatibilizer within the
LDPE matrix and a percolation PANI network within
the compatibilizer. This was based on an idea of
double percolation, described theoretically by Levon
et al. [283] and Knacstedt and Roberts [284]. Yang
et al. [282] suggested that if the above picture is
correct, the resulting percolation threshold might
depend on the length of the alkyl chain in gallic acid
esters, as well as on the nature of the substituents in
the phosphoric acid esters used for protonation
of PANI. This tendency was observed experimen-
tally. Specifically, under identical conditions, the
percolation threshold of the PANI(BEHP)0.5/LDPE–
LG (LDPE:LG ¼ 78:22) composites decreased in
the sequence: propyl gallate . octylgallate . lauryl
gallate. The same was evaluated for different
dopants: three aliphatic esters (di-i-butyl phosphate
(DBP), di-i-octyl phosphate (or bis(2-ethylhexyl)hy-
drogenphosphate-BEHP) and di-i-hexadecyl phos-
phate (DHDP)), and one aromatic ester (DPHP). If
identical processing conditions were used (22 wt%
LG, T ¼ 150 8C, t ¼ 10 min, rotation speed
100 rpm), the percolation threshold decreased in the
following order: DPHP . DBP . DHDP . BEHP.
This suggested that the alkyl chains facilitated the
formation of a continuous percolation network of
PANI in the presence of gallic acid esters. If
unsubstituted aromatic diesters (DPHP) were used
as the dopant for PANI, the gallic acid esters had no
compatibilizing properties. Yang et al. [282] con-
cluded that the dispersion of PANI in LDPE must be
mediated by alkyl chains in the dopant, the
compatibilizer (LG) and the matrix (LDPE).
3.2.3. Temperature effects and ageing of doped PANI
and its composites
The stability of the conductivity and other proper-
ties under operating conditions must be considered for
practical application of PANI blends and composites.
This requires an understanding of different aspects of
thermal action on the materials, and deserves a
separate consideration. Here we present such import-
ant aspects as the effects of temperature on the
properties of PANI and its composites and blends,
particularly including the thermal stability and ageing
of the materials. In part, we have considered these
effects above in connection with processibility,
electronic transport and mechanical properties.
In some cases heating (annealing) samples of
PANI increases their conductivity. It has been
mentioned that in the case of PANI-DBSA, heating
[268] accelerates the doping reaction of PANI by
DBSA, accompanied by a phase transition from a
paste like material to a semi-solid material. Berner
et al. [285] observed an annealing effect after
moderate heating of PANI-CSA films in ambient air
(typically for 30 min at 135 8C), accompanied by an
increase of crystallinity, while the electronic transport
properties improved to a more metallic behavior. The
reflectance spectra of such films aged for some hours
showed distinct evolution stages [286]: (i) increase of
the metallic character after a time of some hours
ageing at 135 8C, (ii) continuous degradation of the
optical conductivity (the real part of the frequency-
dependent complex conductivity) without variation of
the dopant content over a period of 200 h and (iii)
accelerated oxidation and loss of dopant for higher
aging times. Davenas and Rannou [286] concluded
that the existence of a physical ageing stage led to an
improvement of the structural order at the mesoscopic
scale, followed by a stage in which dramatic
alterations of the molecular structure were induced
through chemical degradation.
Amano et al. [166] compared the thermal
stability in air and inert nitrogen conditions of two
polyanilines prepared by aniline polymerization
utilizing two different oxidants, APS and ammonium
dichromate (ADC) in aqueous TSA. They concluded
that the use of ADC allowed the preparation of
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1739
PANI doped with TSA. However, when using APS,
PANI doped with sulfuric acid was synthesized,
despite the presence of TSA. The authors believed
that the dopant (sulfuric acid) originated from the
APS during the oxidation of aniline. The difference
between the samples was manifest in their tempera-
ture dependent conductivity behavior. Thus, in
temperature range from 100 to 180 8C in air and
nitrogen atmosphere, the conductivity of PANI
prepared from APS decreased monotonically with
time. The decreasing conductivity was governed
approximately by first-order kinetics, and its domi-
nant cause was explained by an addition of sulfate
to the PANI aromatic ring. By contrast, the
conductivity behavior of PANI prepared when
using ADC was similar to that discussed above for
PANI-DBSA and PANI-CSA [268,287,288], with
the conductivity increasing with ageing time at
130 8C, regardless of the atmosphere, and showed a
peak with ageing time at 160 8C [166].
Phosphoric acid diesters protonated PANI also
demonstrated a maximum in the dependence of the
conductivity on temperature. Specifically, the maxi-
mum occurred at ,110 8C for PANI(DiOHP)0.3,
,160 8C for PANI(DPHP)0.56 [223]; ,140 8C for
PANI-SPDPAA,,120 8C for PANI-SPDP [280]; and
,120 8C for PANI doped with phosphoric acid
monoesters (3-pentadecylphenilphosphoric acid)
[287]. Mass spectroscopic studies [223] showed that
degradation of the phosphoric acid diesters protonated
PANI was caused by thermal decomposition of the
dopant according to the scheme:
CH3CH2CH2CH2 –O–PðOÞðOHÞ–O–CH2CH2CH2
CH3 ! 2CH2yCHCH2CH3 þ H3PO4
At the same time, Niziol and Laska [288] found
that even at ambient conditions PANI doped with
DiOHP showed time a dependent conductivity during
ageing for a long time. Thus, it increased by about one
order of magnitude during the first year of ageing in
ambient conditions, and then decreased from one to
two orders of magnitude after six years. Similarly, an
increase of conductivity upon two-year ageing was
observed in cellulose acetate blends containing PANI
doped with phenylphosphonic acid [289]. Rannou
et al. [289–291] attributed a decrease of conductivity
of doped PANI at high temperature to its chemical
degradation, caused by three main processes for
the example of PANI-CSA and PANI-HCl: (i)
dedoping, (ii) oxidation/hydrolysis/chain scission,
and (iii) chemical cross-linking (Fig. 11). The first
one is highly dependent on the PANI-protonating
agent, while the other two seem to be general features
of chemical modifications for a thermo-oxidative
ageing of the PANI backbone [289]. Specifically, for
the HCl dopant case, TGA and elemental analysis data
gave evidence of several chemical transformations: (i)
a slight dedoping due to HCl evolution, (ii) an
oxidation of the polymer backbone, and (iii) a
chlorination of the rings. The leading process in
PANI-HCl degradation was found to be the ring
chlorination of PANI rather than HCl evolution,
which account for only 10% of the global phenom-
enon during degradation in air at 140 8C. A more
complex situation was observed for PANI-CSA aging
at 135 8C, in which in situ thermal degradation of
CSA proceeded. The process of PANI-CSA degra-
dation involved: (i) dedoping, (ii) CSA desulfonation,
fragmentation and sulfonation of PANI backbone, (iii)
oxidation, and (iv) chemical cross-linking by for-
mation of interchain tertiary amine bonds [289,290].
Han et al. [292] compared the dependence of the
conductivity on temperature of PANI-DBSA and
PANI-CSA, obtained by dipping EB base films in 1 M
aqueous solution of DBSA and CSA, respectively.
The conductivity of PANI-CSA was higher than
PANI-DBSA, and decreased steeply after about
188 8C for both samples, with that for PANI-CSA
dropping more remarkably after that temperature. At
the same time, whereas the conductivity PANI-DBSA
increased with increasing temperature above 100 8C,that of PANI-CSA decreased slightly, perhaps due to
the evaporation of moisture hydrogen-bonded with
PANI. The authors believed [292] that enhanced
molecular motion of DBSA and PANI with increasing
temperature might be the cause of the conductivity
increment after 100 8C. The higher stability of PANI-
DBSA in comparison with PANI-CSA was due to
higher resistance against deprotonation, and slower
diffusion of DBSA than CSA from PANI on thermal
ageing.
Wang et al. [293] found that treatment of PANI
doped with H2SO4, TSA, or 5-sulphosalicylic acid at
220 8CunderN2 atmosphere for 2 h predominant led to
undoped PANI. The authors found that this process
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531740
was accompanied by a lower quinoid segment content
in the polymer chain. They concluded that a cross-
linking reaction and evolution of the dopant had
occurred during the heat treatment process. This agrees
with a scheme of Rannou et al., see Fig. 11 [289]. The
doped PANI samples displayed no distinctive loss of
conductivity when treated at temperatures 40–200 8C
for 2 h [293]. For temperatures over 200 8C, their
conductivity began to decrease very fast. Specifi-
cally, at 220 8C for 2 h, all conductivities dropped
below 1024 S/cm. Treatment at 220 8C only for
30 min of PANI-DBSA led to a four order loss in
the conductivity, from 120 to 0.01 S/cm, indicated that
220 8Cmight be a ‘dead point’ for sulfonic acid doped
PANI [293].
Tsubakihara et al. [294] also studied the thermal
ageing of the conductivity of PANI-H2SO4 in the
narrower temperature range from 50 to 210 8C with
results differing somewhat from the data ofWang et al.
[293]. Specifically, they found a decrease in
the conductivity at ageing temperatures up to 90 8C,
where the removal of moisture and the reduction of
structural order between polymer chains took place.
The second step of conductivity decrease was found at
Fig. 11. Chemical degradation mechanisms of PANI-HCl and PANI-CSA aged under air [289]. Reproduced from Rannou et al. by permission of
Synth Met 1999;101(1–3):823. q 1999 Elsevier Science Ltd, Oxford, UK.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1741
a temperature higher than 190 8C. Thermally induced
removal of sulfuric acid, and/or some kinds of
chemical reactions should break the formed polaron
band and suppressed the conductivity.
Shacklette et al. [6,80] have found that the
conductivity decay of EB salts varies with time at a
given temperature according to a function of the form:
s ¼ s0 expð2ðt=tÞaÞ; where s is the conductivity at
time t; s0 is the initial conductivity at time t ¼ 0; t isan experimentally determined characteristic decay
time; and a is an experimentally determined par-
ameter for a given sample at each temperature [6,80].
The value of a is typically in the range of 0.77–1.0. A
characteristic half-life of the conductivity can be
determined at each temperature with the help of this
equation, from the value of t and a determined at that
temperature according to the relation: t1=2 ¼ ðln 2Þ1=a;where t1=2 is the time required for the conductivity to
decrease by half. The half-lives followed an Arrhenius
exponential as a function of temperature: t1=2 ¼ ðt1=2Þ0expðEa=kTÞ; where k is the Boltzmann constant. The
authors have determined the important, for practice,
conductivity half-lives for some PANI compositions
(Table 6) [80].
Unlike this, according to Rannou et al. [289]
the kinetics of the conductivity decay of PANI-CSA
films recorded during accelerated ageing tests in air,
performed for seven temperatures between
85 and 175 8C, could be described by classical
Arrhenius low (Fig. 12). The normalized conductivity
loss (s0 ¼ conductivity of unaged film) was
described by two consecutive processes: (i) first one
was an exponential decay, where time was constant, t(h), followed an Arrhenius law (see Eq. (6)),
s=s0 ¼ expð2t=tÞ with t ¼ t0 expðEa=kTÞ ð7Þ
(ii) a second one characterized by a lower rate of
degradation. This description was valid for PANI-
CSA films in the 85 8C , T , 175 8C range and over
2.5 orders of conductivity magnitude. The validity of
the first processes has been used to determine the
activation energy Ea of the global ageing process to
give activation energies of 1.02 and 1.15 eV for films
made with EBI and EBII, respectively (EBI and EBII,
had inherent viscosities of 0.69 and 2.55 dl/g,
respectively, for 0.1 wt% solution in 96 wt%
H2SO4). These results were subsequently used to
predict the long-term behavior of the conductivity.
For PANI-CSA films made with EBI and EBII,
submitted to an isothermal ageing procedure at
50 8C in air, half-life time parameters t1=2 longer to 7
and 34 years were calculated, respectively.
To gain insight into the ageing process, Genoud
et al. [295] used weight uptake data and ESR line
broadening upon oxygen exposure for PANI-CSA
samples after aging at 135 8C for various times. When
the adsorbed gas was paramagnetic oxygen this
resulted in a broadening of the polaron ESR line
proportional to the local conductivity. Gas sorption
experiments, and the kinetics of ESR line broadening
and of dc conductivity confirmed the heterogeneous
structure for PANI-CSA films. Specifically, typically
Table 6
Conductivity half-life of PANI compositions [80]
Composition t1=2(170 8C, h)
t1=2(200 8C, h)
t1=2(230 8C, h)
PANI-TSA 20.6 1.8 0.15
PANI-TSA–DBSA
(TSA:DBSA ¼ 7:3)
20.0 1.5 0.21
PANI-TSA–DBSA(heavy) 10.1 1.3 0.23
PANI-TSA–NDSA 54.3 4.1 0.34
PANI-TSA–NDSA(heavy) 98.6 15.8 1.4
NDSA: naphthalene disulfonic acid.
Fig. 12. Logarithm of the reduced conductivity versus aging time
for air-aged PANI-CSA films [289]. Reproduced from Rannou et al.
by permission of Synth Met 1999;101(1–3):823. q 1999 Elsevier
Science Ltd, Oxford, UK.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–17531742
crystalline highly conducting grains were surrounded
by amorphous less conducting regions. Gas sorption
proceeded via diffusion into the amorphous regions.
Ageing resulted in cross-linking, which slowed down
the gas permeation. In the presence of oxygen the
broadening of the ESR line reflected essentially the
conductivity of the most conducting areas, e.g.
crystalline regions. The latter were less sensitive to
ageing than the amorphous, poorly conducting
regions, which controlled the dc conductivity.
Kuo and Chen [296] characterized the thermo-
stability of the conductivity of PANI doped with
DPHP. They found that the conductivity of PANI-
DPHP powder increased with temperature from 240
toþ140 8C, and decreased with temperature from 140
to 180 8C. This correlated with changes in spin density
of the polymer (Fig. 13).
Naturally, the temperature dependence of the
conductivity of doped PANI was also observed for its
composites. Thus, Ikkala et al. [7] found that blends of
3.2 wt% PANI(DBSA)y (Neste Complex) with HDPE
increased their conductivity followed by the slow
decay with increasing temperature in the range of
70–90 8C.Thermal ageing in various conducting composites
of PANI protonated with hydrochloric acid, and
containing polymers with sulfonic or phosphoryl
groups was investigated by Dalas et al. [297]. They
found that the dc conductivity of the composites for
ageing times from 0 to 300 h decreased at 70 8C in
room atmosphere according to the law s ¼ s0 �expð2ðt=tÞ1=2Þ indicating an inhomogeneous structure
of the granular metal type. It was shown that
composite porosity and the presence of sulfonic or
phosphoryl groups retarded the ageing process. Dalas
et al. [297] attributed thermal degradation of the
composites to a release of HCl from the samples,
which reduced the protonated-conducting phase.
Tsanov and Terlemezyan [298] investigated
the change in the conducting properties of PANI/
poly(ethylene-co-vinylacetate) (CEVA) composite
films as a function of time. They found that the
electrical conductivity of the films with low PANI
content (up to 2.5 wt%) increased by several orders of
magnitude over eight months. This accompanied a
decrease in the average conductivity deviations for
these samples, indicating improvement of conductive
pathways within the insulating CEVA matrix. This
improvement was explained [298] by a change of the
PANI distribution (this was called the apparent
concentration) in the matrix polymer, probably due
to flocculation of the PANI phase, followed by
formation of a continuous conductive network. This
explanation correlates with the phase separation found
during storage of PANI/CEVA films, which leads to
formation of PANI enriched (lower side) and PANI
deficient (upper side) layers with a half order of
magnitude difference in their conductivity [299].
Fig. 13. Temperature dependence of spin density and conductivity of PANI-DPHP [296]. Reproduced from Kuo and Chen by permission of
Synth Met 1999;99(2):163. q 1999 Elsevier Science Ltd, Oxford, UK.
A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 1743
Similar structural and conductive pathways changes
during storage were observed within the blends of
PANI-DBSA with some elastomers (styrene–buta-
diene rubber, nitrile–butadiene rubber or ethylene–
propylene–diene terpolymer) [300].
4. Conclusion
On the whole, the reviewed work testifies both to a
great diversity of PANI containing composites, blends
and methods of their production, as well as to a good
potential for practical use. However, being promising
in a technological sense, this diversity can complicate
the choice of the conducting material for a desired
application. For example, some specific differences
are reported for properties of materials (conductivity,
mechanics, etc.) prepared by different teams under
seemingly similar conditions. Obviously, there is a
problem of taking into account and correspondingly,
maintaining at a constant level, all of the factors
effecting these materials. Specifically, the reviewed
data confirm that the properties of PANI composites
and blends are determined by specific physical–
chemical interactions among their components (PANI
with a dopant, PANI with a host polymer, the dopant
with the host polymer), by the method and conditions
of the material formation, by the quantitative ratio of
the material components, by host polymer precondi-
tions depending on a producer, etc. The situation is
complicated when using plasticizers, which change
the mobility of polymer chains and segments in any
amorphous phase of the material. Finally, these
factors affect the supramolecular structure of a
composite/blend material and the distribution of
PANI in the host matrix. Specifically, an important
role here may be played by the degree of crystallinity
of the material, the size and form of the crystallites,
localization of conducting PANI clusters and PANI
percolation network in the amorphous phase of the
host matrix and by the surface of the crystallites.
Control of these several factors will be necessary for
the production of PANI composites/blends with
predetermined properties. In this connection, when
making a decision on the manufacture of any kind of a
PANI containing composite, a producer has to plan
some research work to adjust and to adopt known data
in the field to fit the particular conditions and
materials relevant to the intended use, or even to
find new results to accommodate the properties
needed in the desired material.
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