Some aspects of preparation methods and properties of polyaniline blends and composites with organic...

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: [email protected] (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


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


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


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.


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


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


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


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


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.


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


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].


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


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.


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


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 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


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.


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


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


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


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


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.


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


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


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].


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.


(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.


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


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.


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.


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


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.


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.


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


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.


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


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


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 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.


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.


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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Some aspects of preparation methods and properties of polyaniline blends and composites with organic...

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