Some aspects of preparation methods and properties ...

Prog. Polym. Sci. 28 (2003) 1701–1753 www.elsevier.com/locate/ppolysci

Some aspects of preparation methods and properties of polyaniline blends and composites with organic polymers Alexander Puda,*, Nikolay Ogurtsova, Alexander Korzhenkob, Galina Shapovala a

Institute of Bioorganic Chemistry and Petrochemistry, Ukrainian Academy of Sciences, 50 Kharkovskoye Shosse, 02160 Kiev, Ukraine b ATOFINA, 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 * Corresponding author. Tel.: þ 380-44-559-70-63; fax: þ380-44-573-25-52. E-mail address: [email protected] (A. Pud). 0079-6700/03/$ - see front matter q 2003 Elsevier Ltd. All rights reserved. doi:10.1016/j.progpolymsci.2003.08.001

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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 following the discovery of solution and melt processibility 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 February 2002 and 2003 in Brussels and Barcelona, respectively. Although a variety of ICPs have been synthesized and investigated, polyaniline, polypyrrole, polythiophene and their derivatives are most often considered, 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 composites 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 polymers, 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 electricity dissipation [35 – 37], heating elements (clothing, 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, conductivity and low cost [2,18,46]. As a consequence, its conducting composites are very close to applications on a large scale for the industrial applications mentioned above [25,32 –45]. Nevertheless, 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 presence 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;


Nomenclature ABS

acrylonitrile – butadiene – styrene copolymer AMPSA 2-acrylamido-2-methyl-1-propanesulphonic 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-dienemonomer) 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

† 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;

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poly-1-caprolactone poly(ethylene oxide) poly(ethylene terephthalate) poly(ethylene terethphalateglycol) poly(methyl methacrylate) poly(m-toluidine) poly(o-methoxyaniline) poly(o-toluidine) 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 polyethylene DSC differential scanning calorimetry DTA differential thermal analysis EMI electromagnetic interference ESR electron spin resonance TGA thermogravimetric analysis XPS X-ray photoelectron spectroscopy

PCL PEO PET PETG PMMA PMT POMA POT PPD-T

† 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 thermoplastic polymer, then molded in a hot press or extruder); Naturally, each of these methods has its own advantages and limitations. Specifically, the synthetic

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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 largescale 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 abovementioned methods, and the results of studies of the resulting PANI composites to elucidate the application 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.

[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 paramagnetic 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]:

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 polymerization, 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 leucoemeraldine ðy ¼ 1Þ; through the EB ðy ¼ 0:5Þ; up to the completely oxidized pernigraniline ðy ¼ 0Þ

ð2Þ Obviously, 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 conformational 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(2vinylpyridine-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 composites [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 sonicating a suspension of the preformed submicronic PANIHCl particles in solutions of the matrix polymers. The blend films exhibited an extremely low percolation threshold ðfp Þ 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

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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 polymerization 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 dodecylsulfate) as the continuous phase and an organic (benzene) solution of the host polymer and aniline as the dispersed phase. The corresponding composites 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

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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 hydroxymethylcellulose. 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 Tg latex component was encapsulated within the high Tg conducting 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(Nvinylpyrrolidone)-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 hydrochloride 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 polymerization 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 composite 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, specifically 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 elastomers 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 developed 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 distribution 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

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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.

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

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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 dispersions (, 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 polymerization 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 polysaccharides [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 network formed during aniline polymerization [89,100].

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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-gPANI in 5 wt% NaOH solution when the molar ratio of aniline/acrylamide (An/AM) in the feed composition 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

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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 comparison 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 semiconductor 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 copolymerization of aniline and its derivatives onto a polymer matrix [108]. Anbarasan et al. [103,104,109] investigated 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 peroxydisulphate 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 conductivity 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 displayed 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 polyethylene, 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 formation 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

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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 crystalline 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 the matrix. 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 polypyromellitimide and PVDF matrices, 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]

Nylon 6 PANI-HCl/ Nylon 6

Heat of fusion, DHf (J/g)

Melting point, Tm (8C)

Degree of crystallinity, Xc (%)

55.28 49.40

217.4 215.5

21.0 18.7

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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 induced much 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 conducting 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 ClO2 4 anion penetration. In this interpretation, the larger ClO2 4 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

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.

spectra of two low frequency relaxation processes, connected with interfacial polarization phenomena. Similar relaxation behavior is observed with composites 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 composites, 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 aniline 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)


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 necessary condition here is penetration (diffusion) of aniline, solvent and electrolyte through the coating to its interface with the anode [31], to create the electrochemical 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 condition, 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 polymerization process is directed from anode to cathode in these electrochemical systems [124], the final composite 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 polymerization of aniline in different polymer matrixes (PU, PC, PMMA, poly-p-phenylene terephthalamide/diphenylether, 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 aniline in porous PVDF or sol – gel silica films covering the anode surface produced micro- and nanocomposites 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 polymerization 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

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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 electrochemical 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 polymerization [129] and varying the matrix. This is probably in agreement with data of Bartlett and Simon [130] on the electrocatalytic properties of electrochemically produced PANI/polyacrylic acid and PANI/poly(vinylsulfonate) films at the anode for NADH (reduced nicotinamide adenine-dinucleotide) oxidation. In comparison 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 transportation of aniline and electrolyte ions through PAN to the working electrode. This suggested that the optimum solution composition favored sufficient

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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 compatible 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 reaction 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 copolymerization 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 sulfonated 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, substituted phosphoric, etc.) and an R – functional group (e.g. alkyl substituted aromatics, long alkyl chains, etc.), imparting dopant compatibility with non-polar or weakly polar organic solvents [2 –5]; Cao et al. [2] called this method to dissolve PANI ‘counter-ion induced processibility’. 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. hexafluoro-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 conductivity 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 preparation 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

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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 interconnected fibrils of diameter around 700 nm, spread throughout the entire surface of the blend. Nevertheless, 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 7:588 could 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 associated with 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, is assigned to degradation of a PANI chain, in agreement 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

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diminished in the blend (0.7, 1.2 and 1.3%, respectively). 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 degradation reactions of the polymer chain and dopant structures, such as cross-linking, loss of conjugation, oxidation, decomposition and other reactions, including 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 temperatures of 130 and 150 8C, respectively, was attributed to dopant loss, degradation reactions, and structural and morphological changes [158]. Specifically, treatment 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 selfsupporting. 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 POMATSA (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 percolation threshold was at least 7 times smaller than predicted by theory. The advantage of POMATSA 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, Gonçalves et al. [170] investigated the suitability of different solvents to prepare PU – POMA blend films by casting.


Specifically, DMF, NMP and m-cresol were compared for this purpose, with DMF selected as the most suitable solvent for two reasons: suppressed deprotonation 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, Gonçalves 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 polyelectrolyte 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 selfassembled layers, prepared through alternate adsorption of the components during 3 min immersion in aqueous solutions [171 – 173]. The complexes demonstrated 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

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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 phenomenon [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 POTHNO3 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 ð10r Þ 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

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the polyaniline chains conformation to facilitate charge transfer. Sevil et al. [180] demonstrated that chlorine substituted PANI had enhanced solubility in comparison with pure PANI. Specifically, they prepared 2chloro-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 dehydrochlorination 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 polyanilines, specifically with sulfonic acid ring-substituted polyaniline (SPAN) and poly(aniline-co-N-propanesulfonic 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 among PVA chains and to a partial miscibility. 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 copolymer, 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 copolymers 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 copolymers 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 nucleophilically attacked aniline at the para position. Copolymers were synthesized starting with diamine polymers of polyethyleneoxide, polypropylene oxide, polydimethylsiloxane and polyacrylonitrile-co-butadiene with molecular weight ðMw Þ 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 ðMw Þ from 30,000 to 140,000. Yamaguchi et al. [187] used a typical amineepoxide reaction when treating LEB with phenylglycidylether (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 corresponding soluble copolymers or their blends. However, 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-2acrylamido-2-methyl-1-propanesulfonic acid (PAMPANI) copolymer with good conductivity was prepared by Yin and Ruckenstein [188]. N-(4Anilinophenyl)methacrylamide (APMA) was synthesized via the catalytic aminolysis reaction (from p-aminodiphenylamine and methylmethacrylate), and poly(AMP-co-APMA) (AMP ¼ 2-acrylamido-2methyl-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(AMPco-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]

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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 mcresol 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 submonolayer 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 crystallinity 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

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conductivity, transparency and other properties. Thus, Yoon et al. [193] found an extremely low percolation threshold ðfp < 0:003Þ when investigating the transport properties of PANI-CSA/ PMMA cast from m-cresol. The electrical conductivity of these blends followed the Mott –Deutscher model for variable-range hopping on fractal network, sðTÞ , exp½2ðT0 =TÞg : The g coefficient increased from 1/4 in pure PANI-CSA (inducing variable-range hopping among exponentially localized 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 temperature, 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 PANICSA/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 micrographs 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 interconnected network affected also thermomechanical properties of PANI-CSA/PMMA and PANIDEHEPSA/PMMA composites prepared from mcresol and dichloroacetic acid, respectively. Thus, whereas the matrix underwent an irreversible flow slightly above its glass-rubber transition temperature, 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

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.

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 DiOHP with PS, cast from m-cresol solutions. Using a fluctuation induced tunneling model, they explained the electrical conductivity variations of the blends in the temperature range between 77 and 300 K by a hopping mechanism between conducting clusters separated by thin insulating 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 PANICSA 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


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 temperature 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 cluster 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

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dissolved in decalin (, 5 wt%) and added to a PANI-CSA0.5 solution (, 1 wt%) in m-cresol, keeping 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 conductivity of 1026 and 0.01 S/cm, for blends containing 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 PANICSA 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 rodlike 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

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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 ultrasonic 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 conditions 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(octadecylacrylate) (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. Specifically, 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 concentrations 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 polymerization, 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 conducting polymer blend of thermoplastic PU (synthesized from polytetramethyleneoxide and 4,4-methylenebis(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 PANIDBSA. 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 components. Thus, they found that the electrical conductivity of PANI-DBSA/SPC composites obtained from chloroform solutions was larger by factor approximately 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 PANIDBSA/SPC composite was about 15 wt%. The conductivity difference suggested that the PANIDBSA 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 distribution 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 nonpolar –CH2 – units in their chains) with PANI doped by CSA, DBSA, or MSA, cast from hexafluoro-2propanol [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 twophase systems: Debye – Bueche (D –B) and inverse power law (IPL). At 3 vol% PANI-CSA0.5 concentration 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

1723

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 concentrations. 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 calorimetry 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 interesting data on the effect of the crystalline structure of the matrix polymer on the formation of electrically conducting networks in conducting blends. Specifically, blends of PANI-CSA with the crystalline and amorphous polyamides Nylon 6 and Selar, respectively, 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 developed conductive pathways in the crystalline host blends, due primarily to the crystallization driven exclusion of the conducting polymer into the interspherultic 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-ethylenedioxythiophene)/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

1724


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 by Wang 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(ethyleneglycol)monomethylether (PEGME, Mw ¼ 350). The DSC curves of the blends with different doped PANI loadings showed a shift of the single endothermic peak (at 67 8C in pure PEO) corresponding to a suppressed melting temperature for the PEO crystallites. 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 considered 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 considered 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 assembling of a conductive pathway during transition of the blend from solution to the solid state [219]. Differences 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 macromolecules. This results in an increase of the intermolecular distance, a corresponding decrease of the intermolecular interaction, and the PANI plasticization. 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 composites 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 approximately 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

1725

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 concentration 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 unwrapping 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, Abraham 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 conductivity of the films was about 0.2 S/cm corresponding, 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.

1726


The strong acid MSA was used by Su et al. [234] as the solvent to prepare blend films of PANI and poly(4vinylpyridine) (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 4sulfophthalic acids improve PANI solution processibility [237 –239]. Polyaniline protonated with these acidic esters was soluble in chloroform, diethylketone, hexafluoro-2-propanol, m-cresol and dichloroacetic 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 diesters 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 mechanical 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 critical 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 macromolecular 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 decomposition 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-co2-acrylamido-2-methyl-1-propanesulfonic acid) (PMMA-co-AMPSA). Emeraldine base, PMMA-coSSA and PS-co-SSA were dissolved in NMP, and the PMMA-co-AMPSA was dissolved in DMF, selected as a better solvent than NMP. 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,

1727

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, hopping and tunneling processes and doping heterogeneities 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 conductivity: (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 corresponding 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

1728


development. Three main approaches exist to produce such materials. The first is realized through the mechanical dispersing of infusible conducting polymers 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 dispersion in thermoplastic matrix polymers and sufficient 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 aggregates 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 ˚ in the skin) is different from for a dopant at or in 50 A ˚ from the the core of these particles (about 50 A 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 surrounded 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 molecules, 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 PMMA matrix. 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 s0 ðw 2 wc Þt ; 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 theoretically predicted value near 2. The values of these parameters for several composites (Table 2) demonstrate 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)

Random filling (3D) PANI in PETG PANI in PCL a

wc (vol%)

s0 a (S/cm)

t

0.15–0.30 0.062 0.046

– 93 292

1.6–2 2.8 1.9

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 PANITSA/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

1729

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 consistent 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 nonconducting 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 metalliclike 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=T 1=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=2 Þ

ð3Þ

where 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 PANITSA/PMMA blends. For these cases the composite expression for conductivity was found to be [254]

s21 ¼ r expð2Tm =TÞ þ t expððT0 =TÞ1=2 Þ

ð4Þ

where 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 at 300 K and fraction m of the resistance at 300 K arising from the first (metallic) term in Eq. (4), given by Ref. [255]: Fig. 4. Themperature dependence of conductivity of two PANITSA/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.

m ¼ r sð300Þexpð2Tm =300Þ

ð5Þ

The value of sð300Þ determined the overall conductivity magnitude, while m indicated the extent of the reduction in slope near room temperature

1730


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 PANI-TSA(33%)/PMMA PANI-TSA(40%)/PETG PANI-TSA(30%)/PETG PANI-TSA(20%)/PETG PANI-TSA(15%)/PETG PANI-TSA

30 13 3.6 0.91 0.10 0.013 18

130 60 770 650 950 1350 1040

0.09 0.11 0.11 0.06 0.03 0.02 0.28

The value of Tm was taken as 2000 K.

(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, showing 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 demonstrated by Zilberman et al. [258,259]. They investigated 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% PANITSA for heterochain polar polymers (PCL, CoPA) and plasticized PS-DOP blends. By contrast, PANITSA blends with non-polar carbochain polymers (LLDPE, LDPE, PS) were conductive only at PANITSA 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 continuous 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). Moreover, 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% PANITSA) only a few very small particles were registered. As a consequence, the higher conductivity at relatively low PANI-TSA content in PCL and CoPA was


1731

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.

assigned to the higher levels of dispersability and structuring of the PANI-TSA particles within the matrix polymer. Thus, the PCL matrix in the PANITSA(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 semicrystalline 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 reaggregation of PANI particles in a matrix polymer. Indeed, interaction of the matrix polymer with the conducting polymer effects dispersion of the conducting 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 interactions 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 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 PCL CoPA LLDPE PS DOP

23.9 17.8 24.2 16.8 19.5 20.9

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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 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 PANITSA. 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-percolation 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 PANITSA 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 Table 4 and Ref. [258]) was found to be similar to that of CoPA, and to be much higher than that of LLDPE, so that only a portion of the PANI particles 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 compatible (PS-DOP) and incompatible (LLDPE) with PANITSA. Hence, one might 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


1733

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.

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 weakening 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. Specifically, 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 polyphenylene 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 PANITSA/Noryl blends had good mechanical properties, with conductivity in a range useful for the production of plastic parts able to dissipate electrostatic electricity. 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 loading in the range of 1.1 –1.5 GPa. Eq. (6) was proposed to correlate the two variables analyzed with the elongation at break ð1b Þ 1b ¼ 9:99 þ 0:097R 2 1:78P

ð6Þ

with 1b given in percentage; R; the rotor speed in rpm; P is the PANI concentration in percentage. From

1734


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 PANITSA 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 dicumylperoxide. 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 PANITSA content, followed by a decrease at higher PANITSA 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

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 ðsb Þ; elongation at break ð1b Þ and gel fraction (GF) [261] PANI concentration (phr)

E (MPa)

sb (MPa)

1b £ 1022 (%)

GF

0 5 10 20 30 40 50

3 3 5 9 14 13 26

5.4 7.6 6.6 6.0 6.5 8.0 7.9

6.0 9.4 6.4 3.5 2.8 2.0 1.0

0.94 0.96 0.96 0.96 0.97 0.99 0.99

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 conducting 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:

Material

PANI(DBSA)y/log s (wt%)

High density polyethylene LDPE Polypropylene PS Impact modified PS PVC Poly(styrene–ethylene/ butylene– styrene)

3.2/24.33; 7.8/22.12 2.4/24.31; 4.3/21.67 1.9/24.33; 7.6/21.77 2.8/26.23; 7.9/22.05 4.0/21.78; 7.2/21.05 1.5/24.37; 2.4/21.22 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 timedependent 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

Fig. 8. Themperature dependence of the storage modulus of PANIDBSA 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.

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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 transition, 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)0.5 complex 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

1736


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 John Wiley & Sons limited.

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 consequence, 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 PANIDBSA(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 PANIDBSA 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 PANIDBSA 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 PANIDBSA(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 crosslinking. However, PANI-DBSA(m)/EPDM and PANI-DBSA(r)/EPDM blends showed partial


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 PANIDBSA(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 agglomerates 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 conditions 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 PANIDBSA 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

1737

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 composites 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 dissipation 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 PANISPDA/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 PANISPDPAA 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

1738


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 Synth Met 2000;114(1):27. q2000 Elsevier Science Ltd, Oxford, UK.

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 Tg with 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 conductivity exceeding 10 S/cm [4]. Polyaniline freestanding 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 conductive blends demonstrated a low percolation threshold (e.g. 6 wt% for PANI(DPHP)0.5/PVCTCP) [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 processing 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]


supposed that LG acted as a compatibilizer, which significantly modified interactions between immiscible 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 solidification 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 microphase 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 doublepercolation 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 experimentally. 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)hydrogenphosphate-BEHP) and di-i-hexadecyl phosphate (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] concluded that the dispersion of PANI in LDPE must be

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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 properties 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 important 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 frequencydependent 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

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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 temperature 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 dominant 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 maximum 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: CH3 CH2 CH2 CH2 – O – PðOÞðOHÞ – O – CH2 CH2 CH2 CH3 ! 2CH2 yCHCH2 CH3 þ H3 PO4 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 phenomenon 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 degradation involved: (i) dedoping, (ii) CSA desulfonation, fragmentation and sulfonation of PANI backbone, (iii) oxidation, and (iv) chemical cross-linking by formation 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 PANIDBSA 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 8C under N2 atmosphere for 2 h predominant led to undoped PANI. The authors found that this process


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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.

was accompanied by a lower quinoid segment content in the polymer chain. They concluded that a crosslinking 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. Specifically, 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 8C might 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 of Wang 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

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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=tÞa Þ; where s is the conductivity at time t; s0 is the initial conductivity at time t ¼ 0; t is an experimentally determined characteristic decay time; and a is an experimentally determined parameter 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=2 Þ0 expðEa =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

Table 6 Conductivity half-life of PANI compositions [80] Composition

t1=2 t1=2 t1=2 (170 8C, h) (200 8C, h) (230 8C, h)

PANI-TSA PANI-TSA–DBSA (TSA:DBSA ¼ 7:3) PANI-TSA–DBSA(heavy) PANI-TSA–NDSA PANI-TSA–NDSA(heavy)

20.6 20.0

1.8 1.5

0.15 0.21

10.1 54.3 98.6

1.3 4.1 15.8

0.23 0.34 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.

(ii) a second one characterized by a lower rate of degradation. This description was valid for PANICSA 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


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 thermostability of the conductivity of PANI doped with DPHP. They found that the conductivity of PANIDPHP powder increased with temperature from 2 40 to þ 140 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

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ageing times from 0 to 300 h decreased at 70 8C in room atmosphere according to the law s ¼ s0 expð2ðt=tÞ1=2 Þ 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.

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Similar structural and conductive pathways changes during storage were observed within the blends of PANI-DBSA with some elastomers (styrene – butadiene rubber, nitrile – butadiene rubber or ethylene– propylene –diene terpolymer) [300].

materials relevant to the intended use, or even to find new results to accommodate the properties needed in the desired material.

References 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 preconditions 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

[1] Shirakawa H, Louis EJ, MacDiarmid AG, Chiang CK, Heeger AJ. Synthesis of electrically conducting organic polymers: halogen derivatives of poly(acetylene), (CH)x. J Chem Soc, Chem Commun 1977;16:578 –9. [2] Cao Y, Smith P, Heeger AJ. Counterion induced processibility of conducting polyaniline and of conducting polyblends of polyaniline in bulk polymers. Synth Met 1992; 48(1):91–7. [3] Heeger AJ. Polyaniline with surfactant counterions-conducting polymer materials which are processible in the conducting form. Synth Met 1993;57(1):3471– 82. ¨ sterholm J-E, Smith P, Heeger AJ, Laska J, [4] Pron A, O Zagorska M. Processable conducting polyaniline. Synth Met 1993;57(1):3520 –5. [5] Cao Y, Smith P, Heeger AJ. Counterion induced processibility of conducting polyaniline. Synth Met 1993;57(1): 3514–9. [6] Shacklette LW, Han CC, Luly MH. Polyaniline blends in thermoplastics. Synth Met 1993;57(1):3532–7. [7] Ikkala OT, Laakso J, Va¨kiparta K, Virtanen E, Ruohonen ¨ sterholm J-E, Cao H, Ja¨rvinen H, Taka T, Passiniemi P, O Y, Andreatta A, Smith P, Heeger AJ. Counterion induced processibility of polyaniline-conducting melt processible polymer blends. Synth Met 1995;69(1 –3): 97 –100. [8] MacDiarmid AG. Polyaniline and polypyrrole: where are we headed? Synth Met 1997;84(1 –3):27–34. [9] Malinauskas A. Chemical deposition of conducting polymers. Polymer 2001;42(9):3957–72. [10] Malinauskas A. Electrocatalysis at conducting polymers. Synth Met 1999;107(2):75–83. [11] Pron A, Rannou P. Processible conjugated polymers: from organic semiconductor to organic metals and superconductors. Prog Polym Sci 2002;27:135– 90. [12] Mortimer RJ. Organic electrochromic materials. Electrochim Acta 1999;44(18):2971–81. [13] Schultze JW, Morgenstern T, Schattka D, Winkels S. Microstructuring of conducting polymers. Electrochim Acta 1999;44(12):1847–64. [14] Bhattacharya A, De AJ. Conducting polymers in solution: progress toward processibility. J Macromol Sci, Rev Macromol Chem Phys 1999;C39(1):17– 56. [15] Novak P, Muller K, Santhanam KSV, Haas O. Electrochemically active polymers for rechargeable batteries. Chem Rev 1997;97(1):207–81. [16] Gospodinova N, Terlemezyan L. Conducting polymers prepared by oxidative polymerization: polyaniline. Prog Polym Sci 1998;23(8):1443 –84.

A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 [17] Bartlett PN, Cooper JM. A review of the immobilization of enzymes in electropolymerized films. J Electroanal Chem 1993;362(1/2):1– 12. [18] Pud AA. Stability and degradation of conducting polymers in electrochemical systems. Synth Met 1994;66(1):1–18. [19] Roncali J. Conjugated poly(thiophenes): synthesis, functionalization, and applications. Chem Rev 1992;92(4): 711–38. [20] Heinze J. Electrochemistry of conducting polymers. Synth Met 1991;43(1/2):2805–23. [21] Deronzier A, Moutet JC. Functionalized polypyrroles. New molecular materials for electrocatalysis and related applications. Acc Chem Res 1989;22(7):249–55. ¨ sterholm J-E. Critical aspects of organic [22] Passiniemi P, O polymer batteries. Synth Met 1987;18(1–3):637–44. [23] Sitaram SP, Stoffer JO, O’Keefe TJ. Application of conducting polymers in corrosion protection. J Coat Technol 1997;69(866):65–71. [24] Huang JC. EMI shielding plastics: a review. Adv Polym Technol 1995;14(2):137 –50. [25] Bhattacharya A, De A. Conducting composites of polypyrrole and polyaniline: a review. Prog Solid State Chem 1996;24(3):141–81. [26] Negi YS, Adhyapak PV. Development in polyaniline conducting polymers. J Macromol Sci, Polym Rev 2002; C42(1):35–53. [27] Niwa O, Tamamura T. Electrochemical polymerization of pyrrole on polymer-coated electrodes. J Chem Soc, Chem Commun 1984;13:817– 8. [28] De Paoli M-A, Waltman RJ, Diaz AF, Bargon J. Conductive composites from poly(vinyl chloride) and polypyrrole. J Chem Soc, Chem Commun 1984;15:1015–6. [29] Han MG, Im SS. Morphological study of conductive polyaniline/polyimide blends. I. Determination of compatibility by small-angle X-ray scattering method. Polymer 2001;42(17):7449–54. [30] Jeevananda T, Siddaramaiah. Synthesis and characterization of polyaniline filled PU/PMMA interpenetrating polymer networks. Eur Polym J 2003;39(3):569–78. [31] Anand J, Palaniappan S, Sathyanarayana DN. Conducting polyaniline blends and composites. Prog Polym Sci 1998; 23(6):993–1018. [32] Taka T. EMI-shielding measurements on poly(3-octyl thiophene) blends. Synth Met 1991;41(3):1177–80. [33] Cottevieille D, Le Me´haute´ A, Challioui C, Mirebeau P, Demay JN. Industrial applications of polyaniline. Synth Met 1999;101(1–3):703– 4. [34] Ma¨kela¨ T, Sten J, Hujanen A, Isotalo H. High-frequency polyaniline shields. Synth Met 1999;101(1 –3):707–7. [35] Jonas F, Schader L. Conductive modifications of polymers with polypyrroles and polythiophenes. Synth Met 1991; 41(1–3):831–6. [36] Jonas F, Heywang G. Technical applications for conductive polymers. Electrochim Acta 1994;39(8/9):1345–7. [37] Stenger-Smith JD. Intrinsically electrically conducting polymers. Synthesis, characterization and their applications. Prog Polym Sci 1998;23(1):57– 79.

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[38] Han C, Shacklette LW. US Patent 5,378,404; 1995. [39] Bakhshi AK. Electrically conducting polymers-from fundamental to applied-research. Bull Mater Sci 1995;18(5): 469 –95. [40] Hanhi K, Lonnberg V, Pyorala K, Loennberg V, Pyoeraelae K. Electroconductive thermoplastic elastomer blend-comprising doped polyaniline blended with ABA styrene block copolymer or olefin polymer blended with crosslinked elastomer. WO Patent 9,706,213; 1997. [41] Misoska V, Ding J, Davey JM, Price WE, Ralph SF, Wallace GG. Polypyrrole membranes containing chelating ligands: synthesis, characterisation and transport studies. Polymer 2001;42:8571–9. [42] Davey JM, Ralph SF, Too CO, Wallace GG, Partridge AC. Electrochemically controlled transport of metal ions across polypyrrole membranes using a flow-through cell. React Funct Polym 2001;49(2):87–98. [43] Wessling B, Posdorfer J. Corrosion prevention with an organic metal (polyaniline): corrosion test-results. Electrochim Acta 1999;44(12):2139–47. [44] Gangopadhyay R, De A. Conducting polymer composites: novel materials for gas sensing. Sens Actuator B 2001;77: 326 –9. [45] Koul S, Chandra R, Dhawan SK. Conducting polyaniline composite: a reusable sensor material for aqueous ammonia. Sens Actuator B 2001;75(3):151–9. [46] Min G. Conducting polymers and their applications in the film industry: polyaniline polyimide blended films. Synth Met 1999;102(1–3):1163–6. [47] Sckotheim TA, Elsenbaumer RL, Reynolds RJ. Handbook of conducting polymers, 2nd ed. New York: Marcel Dekker; 1998. [48] Huang WS, MacDiarmid AG. Optical-properties of polyaniline. Polymer 1993;34(9):1833 –45. [49] Kang ET, Neoh KG, Tan KL. Polyaniline: a polymer with many interesting intrinsic redox states. Prog Polym Sci 1998; 23(2):277–324. [50] Stafstro¨m S, Bredas JL, Epstein AJ, Woo HS, Tanner DB, MacDiarmid AG. Polaron lattice in highly conducting polyaniline: theoretical and optical studies. Phys Rev Lett 1987;59(13):1464–7. [51] Armes SP. Conducting polymer colloids. Curr Opin Colloid Interf Sci 1996;1(2):214–20. [52] Armes SP, Aldissi MJ. Novel colloidal dispersions of polyaniline. J Chem Soc, Chem Commun 1989;2:88– 9. [53] Armes SP, Aldissi MJ, Hawley M, Beery JG, Gottesfeld S. Morphology and structure of conducting polymers. Langmuir 1991;7(7):1447 –52. [54] Gospodinova N, Mokreva P, Terlemezyan L. Stable aqueous dispersions of polyaniline. J Chem Soc, Chem Commun 1992;13:923–4. [55] Eisazadeh H, Spinks G, Wallace GG. Electrodeposition of polyaniline and polyaniline composites from colloidal dispersions. Polym Int 1995;37(2):87 –91. [56] Stejskal J, Kratochvil P, Helmstedt M. Polyaniline dispersions. 5. Poly(vinyl alcohol) and poly(N-vinylpyrrolidone) as steric stabilizers. Langmuir 1996;12(14):3389–92.

1746


[57] Vincent B, Waterson J. Colloidal dispersions of electricallyconducting, spherical polyaniline particles. J Chem Soc, Chem Commun 1990;9:683– 4. [58] Stejskal J, Sˇpirkova M, Riede A, Helmstedt M, Mokreva P, Prokesˇ J. Polyaniline dispersions. 8. The control of particle morphology. Polymer 1999;40(10):2487–92. [59] Chattopadhyay D, Chakraborty M, Mandal BM. Dispersion polymerization of aniline using hydroxypropylcellulose as stabilizer: role of rate of polymerisation. Polym Int 2001; 50(5):538–44. [60] Banerjee PB, Digar ML, Bhattacharyya SN, Mandal BM. Novel polyaniline dispersions using poly(vinyl methyl-ether) stabilizer. Eur Polym J 1994;30(4):499–501. [61] Quadrat O, Stejskal J, Klason C, McQueen DH. The current noise and conductivity behavior of spherical polyaniline particles dispersed in a poly(vinyl alcohol) matrix. J Phys: Condens Matter 1995;7(17):3287– 94. [62] Banerjee P, Mandal BM. Conducting polyaniline nanoparticle blends with extremely low percolation thresholds. Macromolecules 1995;28(11):3940–3. [63] Banerjee P, Mandal BM. Blends of HCl-doped polyaniline nanoparticles and poly(vinyl-chloride) with extremely low percolation-threshold: a morphology study. Synth Met 1995; 74(3):257–61. [64] Beadle P, Armes SP, Gottesfeld S, Mombourquette C, Houlton R, Andrews WD, Agnew SF. Electrically conductive polyaniline copolymer latex composites. Macromolecules 1992;25(9):2526–30. [65] DeArmitt C, Armes SP. Colloidal dispersions of surfactantstabilized polypyrrole particles. Langmuir 1993;9(3):652– 4. [66] Kuramoto N, Genies E. Micellar chemical polymerization of aniline. Synth Met 1995;68(2):191–4. [67] Yang S, Ruckenstein E. Processable conductive composites of polyaniline poly(alkyl methacrylate) prepared via an emulsion method. Synth Met 1993;59(1):1–12. [68] Ruckenstein E, Yang S. An emulsion pathway to electrically conductive polyaniline polystyrene composites. Synth Met 1993;53(3):283–92. [69] Ruckenstein E, Sun Y. Polyaniline-containing electrical conductive composite prepared by 2 inverted emulsion pathways. Synth Met 1995;74(2):107–13. US Patent 5,508,348. [70] Sun Y, Ruckenstein E. Poly(3-methylthiophene)-rubber conductive composite prepared via an inverted emulsion pathway. Synth Met 1995;74(2):145–50. [71] Yassar A, Roncali J, Garnier F. Aqueous suspension of conducting material from polypyrrole-coated submicronic latex particles. Polym Commun 1987;28(4):103–4. [72] Wiersma AE, Vd Steeg LMA, Jongeling TJM. Waterborne core–shell dispersions based on intrinsically conducting polymers for coating applications. Synth Met 1995;71(1– 3): 2269–70. [73] Lascelles SF, Armes SP. Synthesis and characterization of micrometre-sized, polypyrrole-coated polystyrene latexes. J Mater Chem 1997;7(8):1339–47. [74] Lascelles SF, Armes SP, Zhdan PA, Greaves SJ, Brown AM, Watts JF, Leadley SR, Luk SY. Surface characterization of micrometer-sized, polypyrrole-coated polystyrene latexes:

[75]

[76]

[77]

[78]

[79]

[80]

[81]

[82]

[83]

[84]

[85]

[86]

[87]

[88]

[89]

[90]

verification of a core–shell morphology. J Mater Chem 1997; 7(8):1349–55. Barthet C, Armes SP, Chehimi MM, Bilem C, Omastova M. Surface characterization of polyaniline-coated polystyrene latexes. Langmuir 1998;14(18):5032– 8. Barthet C, Armes SP, Lascelles SF, Luk SY, Stanly HME. Synthesis and characterization of micrometer-sized, polyaniline-coated polystyrene latexes. Langmuir 1998;14(8): 2032–41. Maeda S, Cairns DB, Armes SP. New reactive polyelectrolyte stabilizers for polyaniline colloids. Eur Polym J 1997; 33(3):245–53. Perruchot C, Chehimi MM, Delamar M, Lascelles SF, Armes SP. Surface characterization of polypyrrole-coated polystyrene latex by X-ray photoelectron-spectroscopy. Langmuir 1996;12(13):3245–51. Kim BJ, Oh SG, Han MG, Im SS. Preparation of PANIcoated poly(styrene-co-styrene sulfonate) nanoparticles. Polymer 2002;43(1):111–6. Shacklette LW, Miller GG, Elsenbaumer RL, Han CC, Webling BM, Wessling B. Polyaniline compositions having a surface/core dopant arrangement. US Patent 5,281,363, 1994. Han MG, Cho SK, Oh SG, Im SS. Preparation and characterization of polyaniline nanoparticles synthesized from DBSA micellar solution. Synth Met 2002;126(1):53 –60. Xie HQ, Ma YM, Guo JS. Conductive polyaniline-SBS composites from in-situ emulsion polymerization. Polymer 1999;40(1):261 –5. Xie HQ, Ma YM. Change of conductivity of polyaniline/ (styrene–butadiene–styrene) triblock copolymer composites during mechanical deformation. J Appl Polym Sci 2000; 77(10):2156–64. Xie HQ, Ma YM. Preparation of conductive polyaniline/ chlorosulfonated polyethylene composites via in-situ emulsion polymerization and study of their properties. J Appl Polym Sci 2000;76(6):845–50. Jeon BH, Kim S, Choi MH, Chung IJ. Synthesis and characterization of polyaniline–polycarbonate composites prepared by an emulsion polymerization. Synth Met 1999;104(2): 95 –100. Jeevananda T, Siddaramaiah, Annadurai V, Somashekar R. Studies on SLS doped polyaniline and its blend with PC. J Appl Polym Sci 2001;82(2):383–8. Haba Y, Segal E, Narkis M, Titelman GI, Siegmann A. Polyaniline-DBSA/polymer blends prepared via aqueous dispersions. Synth Met 2000;110(3):189– 93. Morita M. Effects of solvent and electrolyte on the electrochromic behavior and degradation of chemically prepared polyaniline – poly(vinyl alcohol) composite. J Polym Sci Polym Phys 1994;32(2):231–42. Gangopadhyay R, De A, Ghosh G. Polyaniline–poly (vinyl alcohol) conducting composite: material with easy processability and novel application potential. Synth Met 2001; 123(1):21–31. Wan MX, Li M, Li JC, Liu ZX. Transparent and conducting coatings of polyaniline composites. Thin Solid Films 1995; 259(2):188–93.

A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 [91] Park YH, Choi SH, Song SK, Miyata S. Synthesis of highly conducting nylon-6 composites and their electrical-properties. J Appl Polym Sci 1992;45(5):843–51. [92] Byun SW, Im SS. Transparent and conducting nylon 6-based composite films prepared by chemical oxidative polymerization. Synth Met 1993;57(1):3501–6. [93] Mandal TK, Mandal BM. Interpenetrating polymer network composites of polypyrrole and poly(methyl acrylate) or poly(styrene-co-butyl acrylate) with low percolation thresholds. Synth Met 1996;80(1):83–9. [94] Zoppi RA, De Paoli M-A. Chemical preparation of conductive elastomeric blends-polypyrrole/EPDM. 2. Utilization of matrices containing cross-linking agents, reinforcement fillers and stabilizers. Polymer 1996;37(10):1999–2009. [95] Wessling B. Polyaniline on the metallic side of the insulatorto-metal transition due to dispersion: the basis for successful nano-technology and industrial applications of organic metals. Synth Met 1999;102(1–3):1396–9. [96] Wessling B. Conducting polymers as organic nanometals. In: Nalva HS, editor. Handbook of nanostructure materials and nanotechnology, vol. 5. San Diego: Academic Press; 2000. p. 501–75. [97] Dutta P, Biswas S, Ghosh G, De A, Chatterjee S. The dc and ac conductivity of polyaniline–polyvinyl alcohol blends. Synth Met 2001;122(2):455–61. [98] Oh SY, Koh HC, Choi JW, Rhee HW, Kim HS. Preparation and properties of electrically conductive polyaniline–polystyrene composites by in-situ polymerization and blending. Polym J 1997;29(5):404–9. [99] Ogura K, Saino T, Nakayama M, Shiigi H. Surfacefunctionalized polyaniline films. J Mater Chem 1997;7(7): 2363–6. [100] Chakraborty M, Mukherjee DC, Mandal BM. Interpenetrating polymer network composites of polypyrrole and poly(vinyl acetate). Synth Met 1999;98(3):193–200. [101] Xiang Q, Xie HQ. Preparation and characterization of alkalisoluble polyacrylamide-g-polyaniline. Eur Polym J 1996; 32(7):865–8. [102] Ghosh P, Chakrabarti A, Siddhanta SK. Studies on stable aqueous polyaniline prepared with the use of polyacrylamide as the water-soluble support polymer. Eur Polym J 1999; 35(5):803–13. [103] Anbarasan R, Jayaseharan J, Sudha M, Devi JL, Nirmala PV, Gopalan AJ. Peroxosalts initiated graft copolymerization of aniline onto rayon fiber: a kinetic approach. Appl Polym Sci 2001;81(2):468–78. [104] Anbarasan R, Vasudevan T, Gopalan A. Peroxosalts initiated graft-copolymerization of aniline onto wool fiber: a comparative kinetic-study. J Mater Sci 2000;35(3): 617–25. [105] Oh KW, Hong KH, Kim SH. Electrically conductive textiles by in-situ polymerization of aniline. J Appl Polym Sci 1999; 74(8):2094– 101. [106] Oh KW, Kim SH, Kim EA. Improved surface characteristics and the conductivity of polyaniline–nylon 6 fabrics by plasma treatment. J Appl Polym Sci 2001;81(3): 684–94.

1747

[107] Oh KW, Hong KH, Kim SJ, Kim SH. Enhancement of polyaniline deposition on the plasma treated nylon 6 fabric. Polym Mater Sci Engng 2000;82:93–4. [108] Tzou K, Gregory V. Kinetic-study of the chemical polymerization of aniline in aqueous-solutions. Synth Met 1992; 47(3):267–7. [109] Anbarasan R, Vasudevan T, Kalaignan GP, Gopalan A. Chemical grafting of aniline and o-toluidine onto poly(ethylene-terephthalate) fiber. J Appl Polym Sci 1999;73(1): 121 –8. [110] Tishchenko GA, Dybal J, Stejskal J, Kudela V, Bleha M, Rosova EY, Elyashevich GK. Electrical resistance and diffusion permeability of microporous polyethylene membranes modified with polypyrrole and polyaniline in solutions of electrolytes. J Membr Sci 2002;196:279 –87. [111] Elyashevich GK, Terlemezyan L, Kuryndin IS, Lavrentyeva VK, Mokreva P, Rosova EY, Sazanov YN. Thermochemical and deformational stability of microporous polyethylene films with polyaniline layer. Thermochim Acta 2001;374(1): 23 –30. [112] Zhang H, Li CJ. Chemical synthesis of transparent and conducting polyaniline poly(ethylene-terephthalate) composite films. Synth Met 1991;44(2):143–6. [113] Chen SA, Fang WG. Electrically conductive polyaniline poly(vinyl alcohol) composite films: physical-properties and morphological structures. Macromolecules 1991;24(6): 1242–8. [114] Byun SW, Im SS. Physical-properties and doping characteristics of polyaniline – nylon-6 composite films. Polymer 1998;39(2):485 –9. [115] Pud AA, Rogalsky SP, Shapoval GS, Korzhenko AA. The polyaniline poly(ethylene-terephthalate) composite. 1. Peculiarities of the matrix aniline redox polymerization. Synth Met 1999;99(2):175–9. [116] Korzhenko AA, Tabellout M, Emery JR, Pud AA, Rogalsky SP, Shapoval GS. Dielectric-relaxation properties of poly(ethylene-terephthalate)–polyaniline composite films. Synth Met 1998;98(2):157–60. [117] Pud AA, Tabellout M, Kassiba A, Korzhenko AA, Rogalsky SP, Shapoval GS, Houze F, Schneegans O, Emery JR. The poly(ethylene terephthalate)/polyaniline composite: AFM, DRS and EPR investigations of some doping effects. J Mater Sci 2001;36(14):3355–63. [118] Nagaoka T, Nakao H, Suyama T, Ogura K. Polyaniline poly(vinyl alcohol) dispersions for controlled ion exchange of organic sulfonates. Analyst 1997;122(11):1399–404. [119] Neoh KG, Tay BK, Kang ET. Oxidation and ion migration during synthesis and degradation of electroactive polymer– nylon 6 composite films. Polymer 2000;41(1):9 –15. [120] Das B, Kar S, Chakraborty S, Chakraborty D, Gangopadhyay S. Synthesis and characterization of polyacrylamide–polyaniline conductive blends. J Appl Polym Sci 1998;69(5): 841 –4. [121] Hatchett DW, Josowicz M, Janata J. Comparison of chemically and electrochemically synthesized polyaniline. J Electrochem Soc 1999;146(12):4535–8.

1748


[122] Fukunada T, Fuse K, Kobayashi K. Metalization of a plastic. Japan Patent 62-139896, 1987. [123] Niwa O, Kakuchi M, Tamamura TJ. Polypyrrole-based conducting polymer allow films: physical properties and film morphology. Macromolecules 1987;20:749 –53. [124] Pud AA, Shapoval GS, Vinogradnyi PG. Electropolymerization of pyrrole in polymer matrices. Russ J Electrochem 2000;36(4):447. [125] Wang HL, Toppare L, Fernandez JE. Conducting polymer blends: polythiophene and polypyrrole blends with polystyrene and poly(bisphenol A carbonate). Macromolecules 1990;23:1053 –9. [126] Neves S, De Torresi CSI, Zoppi RA. Template synthesis of polyaniline: a route to achieve nanocomposites. Synth Met 1999;101(1–3):754– 5. [127] Neves S, Fonseca CP. Determination of fractal dimension of polyaniline composites by SAXS and electrochemical techniques. Electrochem Commun 2001;3(1):36– 43. [128] Neves S, De Paoli M-A. Photoelectrochemistry of polyaniline supported in a microporous cellulose-acetate membrane. Synth Met 1998;96(1):49–54. [129] Pud AA, Fatyeyevax KYu, Shapoval GS. Electrochemical polymerization of aniline in polyamide and poly(vinyl alcohol) matrices. Theor Exp Chem 2002;38(1):33–6. [130] Bartlett PN, Simon E. Poly(aniline) –poly(acrylate) composite films as modified electrodes for the oxidation of NADH. Phys Chem Chem Phys 2000;2(11):2599 –606. [131] Andreev VN. Synthesis and properties of composite polyaniline-nafion films on platinum. Russ J Electrochem 2001; 37(6):605–11. [132] Andreev VN. Synthesis and properties of composite polyaniline–nafion films on glassy carbon. Russ J Electrochem 2001;37(6):612–4. [133] Park YH, Park CR. Preparation of conducting polyacrylonitrile/polyaniline composite films by electrochemical synthesis and their electroactivity. Synth Met 2001;118(1– 3): 187–92. [134] Chen Y, Kang ET, Neoh KG, Wang P, Tan KL. Surface modification of polyaniline film by grafting of poly(ethylene glycol) for reduction in protein adsorption and plateletadhesion. Synth Met 2000;110(1):47– 55. [135] Chen Y, Kang ET, Neoh KG, Tan KL. Covalent immobilization of invertase onto the surface-modified polyaniline from graft-copolymerization with acrylic-acid. Eur Polym J 2000;36(10):2095–103. [136] Ma ZH, Han HS, Tan KL, Kang ET, Neoh KG. Surface graftcopolymerization induced adhesion of polyaniline film to polytetrafluoroethylene film and copper foil. Eur Polym J 1999;35(7):1279–88. [137] Zhao BZ, Neoh KG, Liu FT, Kang ET, Tan KL. N-alkylation of polyaniline with simultaneous surface graft copolymerization for inducing and maintaining a conductive state. Langmuir 2000;16(26):10540– 6. ¨ sterholm H, Ahjopalo L, Takala A, [138] Vikki T, Pietila¨ LO, O Toivo A, Levon K, Passiniemi P, Ikkala O. Molecular recognition solvents for electrically conductive polyaniline. Macromolecules 1996;29(8):2945 –53.

[139] Wei Y, Focke WW, Wnek GE, Ray A, MacDiarmid AG. Synthesis and electrochemistry of alkyl ring-substituted polyanilines. J Phys Chem 1989;93:495–9. [140] Leclerc M, Guay J, Dao LH. Synthesis and characterization of poly(alkylanilines). Macromolecules 1989;22(2):649–53. [141] Macinnes D, Funt BL. Poly-o-methoxyaniline: a new soluble conducting polymer. Synth Met 1988;25(3):235– 42. [142] Yue J, Epstein AJ, MacDiarmid AG. Sulfonic acid ringsubstituted polyaniline, a self-doped conducting polymer. Mol Cryst Liq Cryst 1990;189:255 –61. [143] Yue J, Wang ZH, Cromack KR, Epstein AJ, MacDiarmid AG. Effect of sulfonic-acid group on polyaniline backbone. J Am Chem Soc 1991;113(7):2665–71. [144] Wei XL, Wang YZ, Long SM, Bobeczko C, Epstein AJ. Synthesis and physical-properties of highly sulfonated polyaniline. J Am Chem Soc 1996;118(11):2545–55. [145] Chen SA, Hwang GW. Water-soluble self-acid-doped conducting polyaniline-structure and properties. J Am Chem Soc 1995;117(40):10055–62. [146] Li XG, Huang MR, Zhu LH, Yang YL. Synthesis and air separation of soluble terpolymers from aniline, toluidine, and xylidine. J Appl Polym Sci 2001;82(4):790– 8. [147] Angelopoulos M, Asturias CE, Ermer SP, Ray E, Scherr EM, MacDiarmid AG, Akhtar MA, Kiss Z, Epstein AJ. Polyaniline-solutions, films and oxidation-state. Mol Cryst Liq Cryst 1988;160:151–63. [148] Andreatta A, Cao Y, Chiang JC, Heeger AJ, Smith P. Electrically-conductive fibers of polyaniline spun from solutions in concentrated sulfuric-acid. Synth Met 1988; 26(4):383–9. [149] Cao Y, Smith P, Heeger AJ. Spectroscopic studies of polyaniline in solution and in spin-cast films. Synth Met 1989;32(3):263–81. [150] Harlev E, Gulakhmedova T, Rubinovich I, Aizenshtein G. A new method for the preparation of conductive polyaniline solutions: application to liquid-crystal devices. Adv Mater 1996;8(12):994 –7. [151] Rannou P, Gawlicka A, Berner D, Pron A, Djurado D. Spectroscopic, structural and transport-properties of conductive polyaniline processed from fluorinated alcohols. Macromolecules 1998;31(9):3007–15. [152] Wang ZH, Ray A, MacDiarmid AG, Epstein AJ. Electron localization and charge transport in poly(o-toluidine): a model polyaniline derivative. Phys Rev B 1991;43(5): 4373–84. [153] Dao LH, Leclerc M, Guay J, Chewalier JW. Synthesis and characterization of substituted poly(anilines). Synth Met 1989;29(1):377 –82. [154] Gazotti WA, De Paoli M-A. High yield preparation of a soluble polyaniline derivative. Synth Met 1996;80(3): 263 –9. [155] Raghunathan A, Kahol PK, McCormick BJ. Electron localization studies of alkoxy polyanilines. Synth Met 1999;100(2):205–16. [156] Mattoso LHC, Malmonge LF. Morphology variation as a function of composition for blends of PVDF and a polyaniline derivative. Polymer 1999;40:513–8.

A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 [157] Rocha IS, Mattoso LHC, Malmonge LF, Gregorio R. Effect of low contents of a polyaniline derivative on the crystallization and electrical-properties of blends with PVDF. J Polym Sci, Polym Phys 1999;37(12): 1219–24. [158] Malmonge LF, Mattoso LHC. Thermal-analysis of conductive blends of PVDF and poly(o-methoxyaniline). Polymer 2000;41(23):8387–91. [159] Zulfiqar S, Zulfiqar M, Rizvi M, Munir A, McNeill IC. Study of the thermal-degradation of polychlorotrifluoroethylene, poly(vinylidene fluoride) and copolymers of chlorotrifluoroethylene and vinylidene fluoride. Polym Degrad Stab 1994; 43(3):423–30. [160] Gupta MC, Umare SS. Studies on poly(ortho-methoxyaniline). Macromolecules 1992;25(1):138– 42. [161] Wei Y, Hsueh KF. Thermal-analysis of chemically synthesized polyaniline and effects of thermal aging on conductivity. J Polym Sci, Polym Chem 1989;27(13):4351–63. [162] Javadi HHS, Zuo F, Angelopoulos M, MacDiarmid AG. Mol Cryst Liq Cryst 1988;160:225–33. [163] Rannou P, Nechtschein M. Aging studies on polyanilineconductivity and thermal-stability. Synth Met 1997;84(1– 3): 755–6. [164] Palaniappan S, Narayana BH. Temperature effect on conducting polyaniline salts: thermal and spectral studies. J Polym Sci, Polym Chem 1994;32(13):2431– 6. [165] Chan HSO, Ho PKH, Khor E, Tan MM, Tan KL, Tan BTG, Lim YK. Preparation of polyanilines doped in mixed protonic acids: their characterization by X-ray photoelectron spectroscopy and thermagravimetry. Synth Met 1989;31(1): 95–108. [166] Amano K, Ishikawa H, Kobayashi A, Satoh M, Hasegawa E. Thermal-stability of chemically synthesized polyaniline. Synth Met 1994;62(3):229–32. [167] Malmonge LF, Mattoso LHC. Electroactive blends of poly(vinylidene fluoride) and polyaniline derivatives. Polymer 1995;36(2):245 –9. [168] Angelopoulos M, Ray A, MacDiarmid AG, Epstein AJ. Polyaniline: processability from aqueous solutions and effect of water vapor on conductivity. Synth Met 1987;21(1): 21–30. [169] Wilson A, Gazotti Ir, Roselena F, De Paoli M-A. Thermal and mechanical behavior of a conductive elastomeric blend based on a soluble polyaniline derivative. Eur Polym J 1999; 35(1):35–40. [170] Gonçalves D, Waddon A, Karasz FE, Akcelrud L. Conducting blends of soluble polyurethane and poly(o-methoxyaniline). Synth Met 1995;74(3):197 –9. [171] Paterno LG, Mattoso LHC. Effect of pH on the preparation of self-assembled films of poly(o-ethoxyaniline) and sulfonated lignin. Polymer 2001;42(12):5239–45. [172] Paterno LG, Constantino CJL, Oliveira ON, Mattoso LSH. Self-assembled films of poly(o-ethoxyaniline) complexed with sulfonated lignin. Colloid Surf B 2002;23(4): 257–62. [173] Paterno LG, Mattoso LHC. Influence of different dopants on the adsorption, morphology, and properties of self-assembled

[174]

[175]

[176]

[177]

[178]

[179]

[180]

[181]

[182]

[183]

[184] [185] [186]

[187]

[188]

[189]

1749

films of poly(o-ethoxyaniline). J Appl Polym Sci 2002;83: 1309–16. Anand J, Palaniappan S, Siddaramaiah, Srinivasan P, Sathyanarayana DN. Spectral, thermal and electrical-properties of poly(o-toluidine) and poly(m-toluidine) PVC blends prepared by solution blending. Eur Polym J 1999;35(3):499 –507. Anand J, Palaniappan S, Sathyanarayana DN. Poly(otoluidine)–polystyrene and poly(m-toluidine)– polystyrene blends-spectral, thermal and electrical-properties. Polymer 1998;39(26):6819–25. Anand J, Palaniappan S, Sathyanarayana DN. Solution blending of poly(o-toluidine and m-toluidine) with PMMA in formic-acid medium: spectroscopic, thermal and electrical behavior. Eur Polym J 2000;36(1):157–63. Ahmed SM, Patil RC, Nakayama M, Ogura K. Characterization of picric-acid-doped poly(o-toluidine) and the picricacid-doped poly(o-toluidine)-induced conductive composite of acrylonitrile–butadiene –styrene. Synth Met 2000;114(2): 155 –60. Ahmed SM. Preparation and degradation of highly conducting polyaniline doped with picric acid. Eur Polym J 2002;38: 1151–8. MacDiarmid AG, Epstein AJ. The concept of secondary doping as applied to polyaniline. Synth Met 1994;65: 103 –16. ¨ , Su¨zer S˛. Doping of 2-Cl-PANI/ Sevil UA, Guven O, Birer O PVC films by exposure to UV gamma-rays and e-beams. Synth Met 2000;110(3):175– 9. Chen SA, Hwang GW. Structures and properties of the watersoluble self-acid-doped conducting polymer blends: sulfonicacid ring-substituted polyaniline/poly(vinyl alcohol) and poly(aniline-co-N-propanesulfonic acid aniline)/poly(vinyl alcohol). Polymer 1997;38(13):3333–46. Bartlett PN, Wallace ENK. The oxidation of ascorbate at poly(aniline)–poly(vinylsulfonate) composite coated electrodes. Phys Chem Chem Phys 2001;3(8):1491 –6. Lu F-L, Wudl F, Nowak M, Heeger AJ. Phenyl-capped octaaniline (COA): an excellent model for polyaniline. J Am Chem Soc 1986;108(26):8311–3. Li S, Dong H, Cao Y. Synthesis and characterization of soluble polyaniline. Synth Met 1989;29(1):E329 –36. Li S, Cao Y, Xue Z. Soluble polyaniline. Synth Met 1987; 20(2):141–9. Kinlen PJ, Frushour BG, Ding Y, Menon V. Synthesis and characterization of organically soluble polyaniline and polyaniline block-copolymers. Synth Met 1999;101(3): 758 –61. Yamaguchi I, Yasuda T, Yamamoto T. Ring-opening graft copolymerization of epoxide with polyaniline: chemical properties and lithium ionic conductivity of the copolymer. J Polym Sci, Polym Chem 2001;39(18):3137–43. Yin WS, Ruckenstein E. Water-soluble self-doped conducting polyaniline copolymer. Macromolecules 2000;33(4): 1129–31. Avlyanov JK, Min Y, MacDiarmid AG, Epstein AJ. Polyaniline: conformational-changes induced in solution by variation of solvent and doping level. Synth Met 1995;72(1):65–71.

1750


¨ sterholm H, [190] Ikkala OT, Pietila¨c LO, Passiniemi P, Vikki T, O ¨ sterholm J-E. Processible polyaniline comAhjopalo L, O plexes due to molecular recognition: supramolecular structures based on hydrogen bonding and phenyl stacking. Synth Met 1997;84(1–3):55–8. ¨ sterholm J-E, Monkman AP, [191] Kugler T, Rasmusson JR, O Salaneck WR. Atomic-force microscopy investigations of morphologies in ultrathin polyaniline films. Synth Met 1996; 76(1–3):181–5. [192] Abell L, Adams PN, Monkman AP. Electrical-conductivity enhancement of predoped polyaniline by stretch orientation. Polymer 1996;37(26):5927– 31. [193] Yoon CO, Reghu M, Moses D, Cao Y, Heeger AJ. Transports in blends of conducting polymers. Synth Met 1995;69(1– 3): 255–8. [194] Reghu M, Yoon CO, Yang CY, Moses D, Heeger AJ, Cao Y. Superlocalization of the electronic wave-functions in conductive polymer blends at concentrations near the percolation-threshold. Macromolecules 1993;26(26): 7245–9. [195] Yoon CO, Reghu M, Moses D, Heeger AJ, Cao Y. Electricaltransport in conductive blends of polyaniline in poly(methyl methacrylate). Synth Met 1994;63(1):47–52. [196] Cao Y, Treacy GM, Smith P, Heeger AJ. Solution-cast films of polyaniline: optical-quality transparent electrodes. Appl Phys Lett 1992;60(22):2711–3. [197] Cao Y, Treacy GM, Smith P, Heeger AJ. Optical-quality transparent conductive polyaniline films. Synth Met 1993; 57(1):3526– 31. [198] Yang CY, Cao Y, Smith P, Heeger AJ. Morphology of conductive, solution-processed blends of polyaniline and poly(methyl methacrylate). Synth Met 1993;53(3):293–301. [199] Fraysse J, Plane`s J, Dufresne A, Guermache A. Thermomechanical studies of poly(aniline)/poly(methyl methacrylate) blends: relationship with conducting properties. Macromolecules 2001;34(23):8143–8. [200] Jousseaume V, Morsli M, Bonnet A, Lefrant S. Aging influence on electronic transport-properties of polyaniline blends. Synth Met 1999;101(1–3):813–4. [201] Jousseaume V, Morsli M, Bonnet A, Tesson O, Lefrant S. Electrical properties of polyaniline – polystyrene blends above the percolation threshold. J Appl Polym Sci 1998; 67(7):1205– 8. [202] Jousseaume V, Morsli M, Bonnet A. Aging of electrical conductivity in conducting polymer films based on polyaniline. J Appl Phys 2000;88(2):960 –6. [203] Jousseaume V, Morsli M, Bonnet A. Thermal behavior of polyaniline films and polyaniline–polystyrene blends. J Appl Phys 2002;84:1848– 55. [204] Cao Y, Qiu J, Smith P. Effect of solvents and cosolvents on the processibility or polyaniline. 1. Solubility and conductivity studies. Synth Met 1995;69(1 –3):187–90. [205] Kuramoto N, Teramae K. Preparation of processable polyaniline doped and complexed with anionic surfactant. Polym Adv Technol 1998;9(3):222–6. [206] Valenciano GR, Job AE, Mattoso LHC. Improved conductivity of films of ultra-high-molecular-weight polyethylene

[207]

[208]

[209]

[210]

[211]

[212] [213]

[214]

[215]

[216]

[217]

[218]

[219]

[220]

[221]

[222]

and polyaniline blends prepared from an m-cresol/decaline mixture. Polymer 2000;41(12):4757–60. Zhang Q, Wang X, Geng Y, Chen D, Jing X. Morphology and thermal properties of conductive polyaniline/polyamide composite films. J Polym Sci, Polym Phys 2002;40(22): 2531–8. Andreatta A, Smith P. Processing of conductive polyaniline– Uhmw polyethylene blends from solutions in nonpolarsolvents. Synth Met 1993;55(2/3):1017–22. Leyva ME, Barra GMO, Soares BG. Conductive polyanilineSBS blends prepared in solution. Synth Met 2001;123(3): 443 –9. Kim MS, Levon K. Blend of electroactive complexes of polyaniline and surfactant with alkylated polyacrylate. J Colloid Interf Sci 1997;190(1):17– 36. Koul S, Chandra R, Dhawan SK. Conducting polyaniline composite: a reusable sensor material for aqueous ammonia. Sens Actuator B 2001;75(3):151–9. Brydson JA. Polymeric materials. London: Butterworths; 1989. Ho KS, Hsieh KH, Huang SK, Hsieh TH. Polyurethane-based conducting polymer blends. I. Effect of chain extender. Synth Met 1999;107(1):65–73. Lee W-J, Kim YJ, Kaang S. Electrical-properties of polyaniline/sulfonated polycarbonate blends. Synth Met 2000;113(3):237–43. Hopkins AR, Rasmussen PG, Basheer RA, Annis BK, Wignall GD. Comparison of phase-separation structures in blends of polyaniline/camphor sulfonic-acid and nylon-6 and nylon-12. Synth Met 1998;95(3):179–84. Annis BK, Wignall GD, Hopkins AR, Rasmussen PG, Basheer RA. Structural investigation of polyaniline/nylon 6 blends by small-angle neutron scattering and small-and wideangle X-ray scattering. J Colloid Interf Sci 1997;35(17): 2765–74. Basheer RA, Hopkins AR, Rasmussen PG. Dependence of transition-temperatures and enthalpies of fusion and crystallization on composition in polyaniline/nylon blends. Macromolecules 1999;32(14):4706–12. Hopkins AR, Reynolds JR. Crystallization driven formation of conducting polymer networks in polymer blends. Macromolecules 2000;33(14):5221–6. Wang Y, Wang X, Li J, Zhang H, Mo Z, Jing X, Wang F. Morphological study on water-borne conducting polyaniline–poly(ethylene oxide) blends. J Polym Sci, Polym Phys 2002;40:605–12. Geng YH, Sun ZC, Li J, Jing XB, Wang XH, Wang FS. Water-soluble polyaniline and its blend films prepared by aqueous-solution casting. Polymer 1999;40(20):5723– 7. ¨ sterholm J-E, Smith P. Processable Pron A, Laska J, O conducting polymers obtained via protonation of polyaniline with phosphoric-acid esters. Polymer 1993;34(20): 4235–40. Laska J, Pron A, Lefrant S. Phosphoric-acid diesters protonated polyaniline: preparation, spectroscopic properties, and processability. J Polym Sci, Polym Chem 1995; 33(9):1437–45.

A. Pud et al. / Prog. Polym. Sci. 28 (2003) 1701–1753 [223] Laska J, Pron A, Zagorska M, Lapkowski S, Lefrant S. Thermally processable conducting polyaniline. Synth Met 1995;69(1–3):113– 5. [224] Pron A, Zagorska M, Nicolau Y, Genoud F, Nechtschein M. Highly conductive composites of polyaniline with plasticized cellulose acetate. Synth Met 1997;84(1– 3):89–90. [225] Pron A, Nicolau Y, Genoud F, Nechtshein M. Flexible highly transparent, and conduct polyaniline – cellulose acetate composite films. J Appl Polym Sci 1997;63(8): 971–7. [226] Laska J, Zak K, Pron A. Conducting blends of polyaniline with conventional polymers. Synth Met 1997;84(1– 3): 117–8. [227] Andreatta A, Cao Y, Chiang JC, Heeger AJ, Smith P. Solution processing of polyaniline. Polym Prepr 1989;30(1): 149–50. [228] Andreatta A, Heeger AJ, Smith P. Electrically conductive polyblend fibers of polyaniline and poly-( para-phenylene terephthalamide). Polym Commun 1990;31(7):275–8. [229] Hsu C-H, Vacasegonds P, Epstein AJ. Polyaniline/PPD-T fibers. Synth Met 1991;41(3):1005–8. [230] Hsu C-H, Shih H, Subramoney S, Epstein AJ. High tenacity, high-modulus conducting polyaniline composite fibers. Synth Met 1999;101(1 –3):677–80. [231] Ogurtsov NA, Pokhodenko VD. Effect of polyelectrolyte swelling on conductivity of polyaniline–nylon-6 composite films. Teor Eksp Khimiya 1998;34(1):14–18. [232] Abraham D, Bharathi A, Subramanyam SV. Highly conducting polymer blend films of polyaniline and nylon-6 by cosolvation in an organic-acid. Polymer 1996;37(23): 5295–9. [233] Anand J, Palaniappan S, Sathyanarayana DN. Solution blending of poly(o-toluidine and m-toluidine) with PMMA in formic-acid medium: spectroscopic, thermal and electrical behavior. Eur Polym J 2000;36(1):157–63. [234] Su M-C, Hong J-L. Morphology of conductive polyaniline blended with poly(4-vinyl pyridine). Polymer 2001;42(7): 3297–300. [235] Adams PN, Devasagayam P, Pomfret SJ, Abell L, Monkman AP. A new acid-processing route to polyaniline films which exhibit metallic conductivity and electrical transport strongly dependent upon intrachain molecular dynamics. J Phys Condens Matter 1998;10(37):8293–303. [236] Pomfret SJ, Adams PN, Comfort NP, Monkman AP. Advances in processing routes for conductive polyaniline fibers. Synth Met 1999;101(1–3):724–5. [237] Olinga TE, Fraysse J, Travers JP, Dufresne A, Pron A. Highly conducting and solution-processable polyaniline obtained via protonation with a new sulfonic-acid containing plasticizing functional-groups. Macromolecules 2000;33(6): 2107–13. [238] Dufour B, Rannou P, Fedorko P, Djurado D, Travers JP, Pron A. Effect of plasticizing dopants on spectroscopic properties, supramolecular structure, and electrical transport in metallic polyaniline. Chem Mater 2001;13(11):4032–40. [239] Kulszewicz-Bajer I, Zagorska M, Niziol J, Pron A, Luzny W. Esters of 5-sulfo-I-phthalic acid as new dopants improving

[240]

[241]

[242]

[243]

[244]

[245]

[246]

[247]

[248]

[249]

[250]

[251] [252]

[253]

[254]

[255] [256]

1751

the solution processibility of polyaniline: spectroscopic, structural and transport-properties of the doped polymer. Synth Met 2000;104(2):125– 31. Genies EM, Lapkowski M, Tsintavis C. Preparation, properties and applications of polyaniline. New J Chem 1988;12(4): 181 –96. Wei Y, Tang X, Sun Y. A study of the mechanism of aniline polymerization. J Polym Sci, Polym Chem 1989;27(7): 2385–96. Tzou KT, Gregory RV. Improved solution stability and spinnability of concentrated polyaniline solutions using N,N0 dimethyl propylene urea as the spin bath solvent. Synth Met 1995;69(1–3):109–12. Yin XH, Yoshino K, Hashizume K, Isa I. Preparation and electrical properties of conducting polyaniline–polycarbonate composite films. Jpn J Appl Phys 1997;36(6A): 3537–40. Chen S-A, Lee H-T. Doping behavior of polyaniline film containing 1-methyl-2-pyrrolidone. Synth Met 1993;55(2/3): 1040–4. Afzali A, Buchwalter SL, Buchwalter LP, Hougham G. Reaction of polyaniline with NMP at elevated temperatures. Polymer 1997;38(17):4439–43. Fu YP, Weiss RA. Protonation of polyaniline with lightly sulfonated polystyrene. Macromol Rapid Commun 1996; 17(7):487–92. Hu HL, Cadenas JL, Saniger JM, Nair PK. Electrically conducting polyaniline–poly(acrylic acid) blends. Polym Int 1998;45(3):262 –70. Moon HS, Park JK. Structural effect of polymeric acid dopants on the characteristics of doped polyaniline composites: effect of hydrogen-bonding. J Polym Sci, Polym Chem 1998;36(9):1431 –9. Sixou B, Travers JP, Barthet C, Guglielmi M. Electronic conduction in polyaniline polyethylene oxide and polyaniline–nafion blends: relation to morphology and protonation level. Phys Rev B 1997;56(8):4604–13. Pelster R, Nimtz G, Wessling B. Fully protonated polyaniline: hopping transport on a mesoscopic scale. Phys Rev B 1994;49(18):12718–23. Wessling B. From conductive polymers to organic metals. Chem Innovat 2001;31(1):34–40. Lennartz W, Mietzner T, Nimtz G, Wessling B. Morphological changes in PANI/PMMA-blends during dispersion studied by SAXS. Synth Met 2001;119(1–3): 425 –6. Malliaris A, Turner DT. Influence of particle size on the electrical resistivity of compacted mixtures of polymeric and metallic powder. J Appl Phys 1971;42(2): 614–8. Kaiser AB, Subramaniam CK, Gilberd PW, Wessling B. Electronic transport-properties of conducting polymers and polymer blends. Synth Met 1995;69(1–3):197–200. Subramaniam CK, Kaiser AB, Gilberd PW, Liu CJ, Wessling B. Solid State Commun 1996;97:235–8. Sheng P. Electronic transport in granular metal films. Phil Mag B 1992;65(3):357– 84.

1752


[257] Srinivasan D, Natarajan TS, Rangarajan G, Bhat SV, Wessling B. Electron spin resonance absorption in organic metal polyaniline and its blend with PMMA. Solid State Commun 1999;110(9):503– 8. [258] Zilberman M, Siegmann A, Narkis M. Melt-processed electrically conductive polymer polyaniline blends. J Macromol Sci, Phys 1998;B37(3):301–18. [259] Zilberman M, Siegmann A, Narkis M. Melt-processed, electrically conductive ternary polymer blends containing polyaniline. J Macromol Sci Phys 2000;B39(3): 333–7. [260] Narkis M, Zilberman M, Siegmann A. On the curiosity of electrically conductive melt processed doped-polyaniline/ polymer blends versus carbon-black/polymer compounds. Polym Adv Technol 1997;8(8):525– 8. [261] Mitzakoff S, De Paoli M-A. Blends of polyaniline and engineering plastics. Eur Polym J 1999;35(10):1791–8. [262] Faez R, Gazotti WA, De Paoli M-A. An elastomeric conductor based on polyaniline prepared by mechanical mixing. Polymer 1999;40(20):5497–503. [263] Cao Y, Smith P, Heeger AJ. WO Patent 92/22911; 1992. [264] Pron A, Luzny W, Laska J. Thermally processable polyaniline protonated with diphenyl phosphate: preparation and structural aspects. Synth Met 1996;80(2):191 –3. [265] Ka¨rna¨ T, Laakso J, Levon K, Savolainen E. US Patent 5,346, 649; 1994. [266] Jussila M, Karhu E, Karna T, Kirmanen P, Laakso J, Suominen M, Kaernae T. WO Patent 9,528,716; 1995. [267] Levon K, Ho KH, Zheng WY, Laakso J, Karna T, Taka T, ¨ sterholm J-E. Thermal doping of polyaniline with dodeO cylbenzene sulfonic-acid without auxiliary solvents. Polymer 1995;36(14):2733–8. [268] Ahlskog M, Isotalo H, Ikkala O, Laakso J, Stubb H, ¨ sterholm J-E. Heat-induced transition to the conducting O state in polyaniline dodecylbenzenesulfonic acid complex. Synth Met 1995;69(1 –3):213–4. [269] Ikkala OT, Lindholm TM, Ruohonen H, Selantaus M, Va¨kiparta K. Phase-behavior of polyaniline dodecyl benzene sulfonic-acid mixture. Synth Met 1995;69(1– 3): 135– 6. [270] Vikki T, Ikkala OT. On the dynamic-mechanical relaxations of polyaniline(dodecyl benzene sulfonic acid)-salt. Synth Met 1995;69(1–3):235–6. [271] Titelman GI, Zilberman M, Siegmann A, Haba Y, Narkis M. Thermal dynamic processing of polyaniline with dodecylbenzene sulfonic acid. J Appl Polym Sci 1997;66(12): 2199–208. [272] Zilberman M, Titelman GI, Siegmann A, Haba Y, Narkis M, Alperstein D. Conductive blends of thermally dodecylbenzene sulfonic acid-doped polyaniline with thermoplastic polymers. J Appl Polym Sci 1997;66(2):243–53. [273] Faez R, De Paoli M-A. A conductive rubber based on EPDM and polyaniline. I. Doping method effect. Eur Polym J 2001; 37(6):1139– 43. [274] Faez R, De Paoli M-A. Elastic polyaniline with EPDM and dodecylbenzenesulfonic acid as plasticizers. J Appl Polym Sci 2001;82(7):1768–75.

[275] Faez R, Martin IM, De Paoli M-A, Rezende MC. Influence of processing time and composition in the microwave absorption of EPDM/PANI blends. J Appl Polym Sci 2002;83(7): 1568–75. [276] Faez R, Schuster RH, De Paoli M-A. A conductive elastomer based on EPDM and polyaniline. II. Effect of crosslinking method. Eur Polym J 2002;38:2459– 63. [277] Davies SJ, Ryan TG, Wilde CJ, Beyer G. Processable forms of conductive polyaniline. Synth Met 1995;69(1–3): 209 –10. [278] Leyva ME, Bara GMO, Gorelova MM, Soares BG, Sens M. Conducting SBS block copolymer–polyaniline blends prepared by mechanical mixing. J Appl Polym Sci 2001;80(4): 626 –33. [279] Koul S, Chandra R, Dhawan SK. Conducting polyaniline composite for ESD and EMI at 101 GHz. Polymer 2000; 41(26):9305–10. [280] Paul RK, Pillai CKS. Melt/solution processable conducting polyaniline with novel sulfonic-acid dopants and its thermoplastic blends. Synth Met 2000;114(1):27–35. [281] Ikkala OT, Pietila¨ LO, Passiniemi P, Cao Y, Andreatta A. European Patent 0,643,397; 1995. [282] Yang JP, Rannou P, Planes J, Pron A, Nechtschein M. Preparation of low density polyethylene-based polyaniline conducting polymer composites with low percolation threshold via extrusion. Synth Met 1998;93(3):169–73. [283] Levon K, Margolina A, Patashinsky AZ. Multiple percolation in conducting polymer blends. Macromolecules 1993; 26(15):4061–3. [284] Knacstedt MA, Roberts AP. Morphology and macroscopic properties of conducting polymer blends. Macromolecules 1996;29(4):1369 –71. [285] Berner D, Davenas J, Djurado D, Nechtschein M, Rannou P, Travers JP. Annealing effect in polyanilineCSA upon moderate heating. Synth Met 1999;101(1–3): 727 –8. [286] Davenas J, Rannou P. Optical reflectance and RBS studies of thermally aged CSA-protonated polyaniline films. Synth Met 1999;102(1–3):1283–4. [287] Paul RK, Vijayanathan V, Pillai CKS. Melt solution processable conducting polyaniline: doping studies with a novel phosphoric-acid ester. Synth Met 1999;104(1–3): 189–95. [288] Niziol J, Laska J. Conductivity of blends of polyaniline with PMMA and cellulose-acetate-aging studies. Synth Met 1999; 101(1– 3):720–1. [289] Rannou P, Nechtschein M, Travers JP, Berner D, Wolter A, Djurado D. Aging of PANI: chemical, structural and transport consequences. Synth Met 1999;101(1–3): 734 –7. [290] Rannou P, Rouchon D, Nicolau YF, Nechtschein M, Ermolieff A. Chemical degradation of aged CSA-protonated PANI films analyzed by XPS. Synth Met 1999;101(1–3): 823–4. [291] Wolter A, Rannou P, Travers JP, Gilles B, Djurado D. Model for aging in HCl-protonated polyaniline: structure,


[292]

[293]

[294]

[295]

[296]

conductivity, and composition studies. Phys Rev B 1998; 58(12):7637– 47. Han MG, Lee YJ, Byun SW, Im SS. Physical properties and thermal transition of polyaniline film. Synth Met 2001; 124(2/3):337–43. Wang XH, Geng YH, Wang LX, Jing XB, Wang FS. Thermal behaviors of doped polyaniline. Synth Met 1995;69(1– 3): 265–6. Tsubakihara H, Miyazaki S, Fujisawa S. Electrical conductivity of polyaniline with sulfuric acid doping. Kobunshi Ronbunshu 1997;54(7):455– 9. Genoud F, Nechtschein M, Rannou P, Escoubes M, Rouyer V. Gas sorption and ESR studies of aging in PANI-CSA. Synth Met 1999;101(1 –3):798–9. Kuo C-T, Chen C-H. Characterization of polyaniline doped with diphenyl phosphate. Synth Met 1999;99(2):163 –7.

1753

[297] Dalas E, Sakkopoulos S, Vitoratos E. Thermal-degradation of the electrical-conductivity in polyaniline and polypyrrole composites. Synth Met 2000;114(3):365–8. [298] Tsanov T, Terlemezyan L. New evidence supporting the inter-facial model of electrically conductive polymer composites. I. Conductivity behavior of polyaniline poly(ethylene-co-vinyl acetate) composites. Polym Polym Compos 1997;5(7):483– 92. [299] Tsanov T, Terlemezyan L. New evidence supporting the interfacial model of electrically conductive polymer composites. 2. Indication of phase separation during ageing of polyaniline – poly(ethylene-co-vinylacetate) composites. Polym Polym Compos 1998;6(1):39–45. [300] Tsanov T, Ditcheva-Kortchakova M, Terlemezyan L. Electrically conductive elastomer composites containing polyaniline. Polym Polym Compos 2000;8(2):115–21.

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