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Phthalonitrile-containing poly(amide imide)s With nanoactuation properties.


Aromatic polyimides are well known because of their excellent thermal stability and electrical insulation properties. They are mainly used in aerospace and electronic industries in the form of films and moldings. However, their applications are limited because rigid polyimides are insoluble and infusible in their imidized form, which leads to process difficulties. The most used method for the synthesis of aromatic polyimides via soluble poly(amic acid) precursors has disadvantages connected with the low stability of the poly(amic acid) solutions and incomplete cyclization of the poly(amic acid)s to polyimide structures (1), (2). To over-overcome these drawbacks, modifications of the polyimide structure are often used, for example, introducing flexible linkages, nonsymmetrical structures, or bulky substituents into the polymer backbones (3-5) Aromatic poly(amide imide)s have been developed as alternative materials offering a compromise between excellent thermal stability and processability. They possess desirable characteristics of the polyamides and polyimides such as high thermal stability and good mechanical properties as well as easy processability. Therefore, a wide range of aromatic poly(amide imide)s have been investigated as high-performance materials for applications, such as moldings, wire-coating enamels, and high-temperature adhesives (6-9).

In recent years, fluorinated functional groups have been introduced into polymer backbones to give polymer materials with improved characteristics. The flexible hexafluor-oisopropylidene (6F) groups introduced between the aromatic and heteroaromatic rings enhance the polymer solubility without sacrifying thermal stability. The retention of high thermal stability is attributed to the strong C--F bonds. Other effects of the 6F groups are the increased glass transition temperature and flame resistance with a concomitant decrease of crystallinity and water absorption. The bulky 6F groups also serve to increase the free volume of the polymers, thus improving its electrical insulating characteristics (10-14). Fluorinated poly(amide imide)s show outstanding capability to be processed into flexible, transparent films, having low dielectric constant and tough mechanical characteristics. The polymer films or coatings can be used in advanced microelectronic and optoelectronic applications (15).

It was proved that the aromatic polyimides having nitrite groups are the most promising materials with piezoelectric properties for use at high temperature (16), (17). The large dipole moment of nitrite group (4.18 D) provides a strong interaction with the applied electric field (18). Moreover, the introduction of nitrile groups into the polymer chains maintains a high thermal stability. Nitrile substituents possess strong bond dissociation energy (>500 kJ [mo1.sub.-1]), which keeps the excellent thermal stability of the polymer (19). The pendant nitrile group on aromatic ring appears to promote adhesion of the polymer to many substrates (20), possibly through polar interaction with other functional groups; it also serves as a potential site for polymer crosslinking (21), (22).

Phthalonitrile monomers or oligomers with different structures have been synthesized and used for the preparation of phthalonitrile resins (23-25). Polymerization of these monomers occurs through nitrite groups by an addition cure mechanism. During this reaction, little or no volatiles are evolved leading to void-free crosslinked product such as triazine, phthalocyanine, and isoindoline (26). The formation of heterocyclic crosslinked structures increases the thermo-oxidative stability and improves mechanical properties. Thus, a soluble aromatic polyamide with bulky pendant phthalonitrile units was prepared from a phthalonitrile-containing diamine. 4-(4-(3,5-diaminoben-zoyl)phenoxy)phthalonitrile, and isophthaloyl chloride. The solvent resistance and thermal properties of the polymer were enhanced by thermal treatment around the glass transition temperature for an extended period of time because of the crosslinking site formation through the thermal polymerization of nitrile groups on pendant phthalonitrile units (26).

In the light of these facts, we synthesized and characterized new fluorinated poly(amide imide)s containing pendant phthalonitrile groups starting from 4,4'-diamino-4"-(3,4-dicyanophenoxy)triphenylmethane, 1, 1,3-bis(4-aminophenoxy)benzene, and a diacid chloride containing a 6F group and imide rings. The properties of these polymers such as solubility, thermal stability, and electrical characteristics were investigated with respect to their chemical structure. The nanometric displacements of polymer films were measured when an electric field was applied on their surface.



4,4'-Diam ino-4"-hydroxy-triphenylmethane (DHTM) was obtained by the reaction of 4-hydroxybenzaldehyde and aniline in the presence of an acid catalyst (aniline hydrochloride), as was previously reported (27). Diamine I was prepared by nucleophilic displacement reaction of 4-nitrophthalonitrile with an aromatic diamine containing a hydroxyl group DHTM, in the presence of anhydrous potassium carbonate, as shown in Scheme 1 (28). Mp: 78.4[degrees]C

(DSC, 10[degrees]C [min.sup.-1]) FTIR (KBr, [cm.sup.-1]): 3477, 3377 and 3220 [cm.sup.-1] (NH-amine), 3030 (C--H), 2230 (CN), 1240 (aromatic ether). [.sup.1]H NMR (DMSO-[d.sub.6], ppm. [delta]): 8.08 (d, 1H), 7.79 (s, 1H), 7.35 (d, 1H), 7.17 (d, 2H), 7.09 (d, 2H), 6.76 (d, 4H), 6.49 (d, 4H), 5.22 (s, CH, 1H), 4.93 (s, [NH.sub.2], 4H).

2,2-Bis[N-(4-chloroformylphenyl)phthalimidyl]hexafluoro-propane, 2, was obtained by treating with thionyl chloride the corresponding dicarboxylic acid resulting from the condensation reaction of 4,4'-(hexafluoroisopropylidene)-diphthalic anhydride with p-aminobenzoic acid, in glacial acetic acid as a solvent and dehydrating reagent, at reflux temperature (29). Mp: 311-313[degrees]C. FT1R (KBr, [cm.sup.-1]): 1780 (C=O of imide ring and COCI), 1720 (C=O of imide ring), 1600 (aromatic), 1390 (C--N), 1210 and 1180 ([CF.sub.3]), 1100 and 720 (imide ring). [.sup.1]H NMR (DMSO-[d.sub.6], ppm, [delta]): 8.22 (2H, d), 8.11 (4H, d), 7.91 (2H, d), 7.73 (2H, s), 7.61 (4H, d).

4-Hydroxybenzaldehyde, aniline, aniline hydrochloride, 4-nitrophthalonitrile, thionyl chloride, 4,4'-(hexafluoroiso-propylidene)diphthalic anhydride, p-aminobenzoic acid, and 1,3-bis(4-aminophenoxy)benzene, 3, were provided by Aldrich and used as received.

Synthesis of Poly(cnnicle intide)s 4

The synthetic route of the polymers 4 is depicted in Scheme 2. The experimental details are described below using polymer 4a as an example. Diamine 1 (0.707 g, 1.7 mmol), N-methyl-2-pyrrolidone (NMP) (13 mL), and pyridine (0.4 mL) were placed in a 100-mL three-necked flask equipped with a mechanical stirrer and nitrogen gas inlet and outlet. The solution was cooled to -1O[degrees]C and diacid chloride 2 (1.222 g, 1.7 mmol) was added with rapid stirring. The content of the flask was kept below 0[degrees]C for 15 min. The cooling bath was then removed and the reaction mixture was allowed to reach room temperature, after which it was stirred for further 8 h. Half of the resulting viscous solution was cast onto glass plates and after evaporating the solvent at 80, 120, 140, 180, and 220[degrees]C, each for 1 11, flexible films were obtained, which were stripped off the plate by immersion in water. The films were dried at 100[degrees]C for 12 11 and used for different measurements. The other half of the polymer solution was diluted by addition of more NMP and the polymer was precipitated by pouring into water. The precipitated product was filtered, washed twice with ethanol under stirring, and dried under vacuum at 100[degrees]C for 12 h.

Polymer 4c was synthesized according to the above procedure, using 1,3-bis(4-aminophenoxy)benzene (0.496 g, 1.7 mmol) instead of 1, whereas copolyamide 4b was prepared starting from 1 (0.353 g, 0.85 mmol), diacid chloride 2 (1.222 g, 1.7 mmol), and 1,3-bis(4-aminophenoxy)benzene (0.248 g, 0.85 mmol).

Polymer 4a. FTIR (KBr, [cm.sup.-1]): 3350, 3060, 2930, 2860, 2232, 1785, 1720, 1665, 1600, 1508, 1368, 1254, 1209, 1188, 719; IH NMR (DMSO-[d.sub.6], ppm, [delta]): 10.35 (s, CONH, 2H), 8.22 (d, 2H), 8.09 (s, 1H), 8.07 (d, 4H), 8.00 (d, IH), 7.82 (s, 1H), 7.77 (d, 6H), 7.62 (d, 4H), 7.40 (d, IH), 7.27 (d, 2H), 7.17 (d, 6H), 5.65 (s, IH).

Polymer 4h. FTIR (KBr, [cm.sup.-1]): 3347, 3075, 2937, 2232, 1786, 1721, 1605, 1505, 1366, 1250, 1207, 717; IH NMR (DMSO-[d.sub.6], ppm, [delta]): 10.39 (s, CONH, 4H), 8.22 (d, 4H), 8.09 (d, 8H), 8.00 (m, 5H), 7.84 (d, 5H), 7.78 (d, 8H), 7.63 (d, 8H), 7.40 (d, 1H), 7.35 (t, 1H), 7.27 (d, 2H), 7.18 (d, 6H), 7.11 (d, 4H), 6.71 (d, 2H), 6.60 (s, 1H), 5.65 (s, 1H).

Polymer 4c. FTIR (KBr, [cm.sup.-1]): 3336, 3098, 1786, 1721, 1665, 1606, 1506, 1368, 1255, 1208, 718; [.sup.1]H NMR (DMSO-[d.sub.6], ppm, [delta]): 10.39 (s, CONH, 2H), 8.22 (d, 2H), 8.08 (d, 4H), 7.99 (s, 1H), 7.83 (m, 6H), 7.63 (d, 4H), 7.35 (t, 1H), 7.11 (d, 4H), 6.71 (d, 2H), 6.60 (s, 1H).


Melting points of the monomers and intermediates were measured on a Melt-Temp II (Laboratory Devices).

The inherent viscosities of the polymers were determined at 20[degrees]C for polymer solutions (0.5 g [dL.sup.-1]) in NMP, using an Ubbelohde viscometer.

FTIR spectra were recorded on a Bruker Vertex 70 at frequencies ranging from 400 to 4000 [cm.sup.-1] using KBr pellets or films.

[.sup.1]H NMR (400 MHz) spectra were performed at room temperature on a Bruker Avance DRX 400 spectrometer, using DMSO-[d.sub.6] as solvent.

The molecular weights and their distributions were determined by gel permeation chromatography (GPC) with a PL-EMD 950 evaporative light scattering detector instrument. Two poly(styrene-codivinylbenzene) gel columns (PLgel 5 [mu]m Mixed-D and PLgel 5 [mu]m Mixed-C) were used as stationary phase, while N,N-dimethylformamide (DMF) was used as the mobile phase. The eluent flow rate was 1.0 mL [min.sup.-1]. Polystyrene standards of known molecular weight were used for calibration.

The water absorption values of the films were determined at 20[degrees]C. The films were vacuum dried at 100[degrees]C for 24 h before testing of the water absorption, for which they were weighted and immersed in deionized water at room temperature for 24 h. The wet films were wiped dry and weighted again. The water absorption of the films was calculated in weight percent as follows:

Water absorption(%) = [([W.sub.2]- [W.sub.1])/[W.sub.1] x 100, (1)

where [W.sub.2] is the weight of film after water absorption and WI is the weight of dry film.

For each him, the water absorption represents the average of three measurements.

The wide-angle X-ray diffraction (WAXD) experiments were performed on a D8 Advance Bruker AXS diffractometer using a CuKu source with an emission current of 36 mA and a voltage of 30 kV. Scans were collected over the 2[theta] = 2[degrees]-40[degrees] range using a step size of 0.01[degrees] and a count time of 0.5 s per step.

Thermogravimetric analysis (TGA) was performed in air on a MOM derivatograph (Hungary) at a heating rate of 10[degrees]C [min.sup.-1]. The initial decomposition temperature is characterized as the temperature at which the sample achieves a 5% weight loss. The temperature of 10% weight loss ([T.sub.10] was also recorded. The maximum decomposition rate temperature, which is the maximum signal in differential thermogravimetry curves, was also recorded.

The dielectric measurements were carried out using a Novocontrol Dielectric Spectrometer (GmbH Germany), CONCEPT 40. The samples were prepared in the form of films with thickness of 20-40 [micro]m. They were sandwiched between two copper electrodes of diameter 20 mm and placed inside temperature-controlled sample cell. The complex permittivity, [epsilon]*(f) = [epsilon]'(f) - i[epsilon]"(f), has been determined in the frequency (f) range from [10.sup.-1] to [10.sup.6] Hz. Temperature was controlled using a nitrogen gas cryostat, and the temperature stability of the sample was better than 0.1[degrees]C. The samples were measured first in the temperature range of -150 to 200[degrees]C (first scan). Then, keeping them into the cell, they were measured again from -150 to 280[degrees]C (second scan). The complex permittivity was converted to the complex dielectric modulus M*(f) according to an equation described in the literature (30). The real (0) and imaginary (M") parts of the dielectric modulus can be calculated from [epsilon]' and [epsilon]":

M' = [epsilon]' / 2 ([[epsilon]'.sup.2] + [[epsilon]".sup.2]) (2)

M' = [epsilon]''/ 2 ([[epsilon]'.sup.2] + [[epsilon]".sup.2]) (2)

The main advantage of this formulation is that the space charge often does not mask the features of the spectra, because of the suppression of high capacitance phenomena in M" plot.

The linear nanodisplacement of polymer films was determined at room temperature, using an experimental setup for linear measurements, based on a Michelson type interferometer using AGILENT 5529A system.


Aromatic poly(amide imide)s 4 were synthesized by low-temperature solution polycondensation of a fluorinated diacid chloride 2 with a mixture of diamine 1 and an aromatic diamine containing flexible ether linkages 3, in various ratios, using NMP as solvent and pyridine as acid acceptor (Scheme 2).

The structure of the resulting polymers was investigated by FTIR and Ill NMR spectroscopy. In the FTIR spectra, the absorption bands that appeared at 3400-3300 and 1665 [cm.sup.-1] were attributed to the NH and carbonyl stretching vibrations of amide groups. All the polymers exhibited strong bands at 1785 and 1720 [cm.sup.-1], which are commonly attributed to the asymmetrical and symmetrical stretching vibrations of carbonyl groups of imide rings; absorption band at 1368 [cm.sup.-1] was due to C--N stretching of imide rings, and absorption band at 719 [cm.sup.-1] was due to imide ring deformation. In FT1R spectra of polymers 4a and 4b, the band appearing at 2232 [cm.sup.-1] was typically for nitrile group of phthalonitrile units coming from the segments of diamine 1. All the polymers exhibited absorption bands at 1188 and 1209 [cm.sup.-1] due to 6F groups and at 1254 [cm.sup.-1] due to aromatic ether linkages (14). Characteristic bands at 3060 and 1600 [cm.sup.-1] were attributed to aromatic C--H stretching and aromatic C=C stretching, respectively. Figure 1 presents the FTIR spectra of polymers 4. Figure 2 shows [.sup.]H NMR spectrum of polymer 4a with the assignment for all the protons. The peaks corresponding to the aromatic protons were situated in the region of 8.22-7.17 ppm. The protons [H.sub.12] and [H.sub.13] of the dianhydride segments appeared at the farthest downfield region of the interval corresponding to aromatic protons. The proton of aliphatic group CH appeared at lower ppm value (5.65 ppm), and the proton of amide group appeared at the highest ppm value (10.35 ppm). From 11-1 NMR spectra of 4, it was found that the composition of the polymers was similar with the composition of the reactants used in synthesis.

The inherent viscosity of polymers was in the range of 0.48-0.62 dL [g.sup.-1]. The molecular weight was determined by GPC. The values of weight-average molecular weight ([M.sub.w]) and number-average molecular weight ([M.sub.n]) were in the domain of 23,250-36,500 g [mo1.sup.-1] and 35,100-51,700 g [mol.sup.-1], respectively. The polydispersity Mw/Mi, was in the interval of 1.38-1.51 (Table 1). The GPC curves showed narrow molecular weight distribution and low quantity of oligomers.

TABLE 1. Properties of polymers 4.

Polymer    [[eta].sub  [M.sub.n]    [M.sub.w](g  [M.sub.w]/       Wuler
         [inh.sup.a]]    (g[mol.  [mol.sup.-1])    [Msub.n]  absorption
              (dL [g.   sup.-1])                                    (%)

4a               0.48     23,250         35,100        1.51        3.48

4b               0.62     36,500         51,700        1.41        3.11

4c               0.57     27,500         38,000        1.38        2.91

(a.) Measured at a concentration of 0.5 g polymer in
100 mL of NMP, at 20[degrees]C.

The polymers were soluble in polar aprotic solvents such as NMP, DMF, or N,N-dimethylacetamide (DMAc), and even in less polar liquids such as pyridine. The good solubility of these polymers can be explained mainly by the presence of voluminous 6F groups, which prevent a strong packing of the chains and thus facilitate the diffusion of small molecules of solvent. The polymers 4 were heated at 280[degrees]C for 2 h. Polymers 4a and 4b exhibited, after this thermal treatment, enhanced solvent resistance; they were insoluble in organic solvents (NMP, DMF, and DMAc) because of the formations of crosslinking sites through the nitrile cure reactions (26).

The water absorption data of films are listed in Table I. It shows that the water absorption values of polymers 4a and 4b were higher than that of polymer 4c not containing phthalonitrile groups. This is attributed to the increase of hydrophilicity of polymers with the increase of polar phthalonitrile group content.

All the films obtained by solution casting and thermal treatment were flexible and tough. The color of polymer film 4c was light yellow. An increase of the color intensity appeared by increasing the phthalonitrile content, probably because of an increase of intermolecular interactions.

The crystallinity of the polymers was measured by WAXD scans. Typical diffraction patterns are presented in Fig. 3. The polymers showed completely amorphous pattern. The amorphous nature of these polymers was reflected in their excellent solubility.

The thermo-oxidative stability was studied by TGA, effectuated in air atmosphere. The main thermal parameters are summarized in Table 2. The polymers began to decompose in the range of 427-450[degrees]C, as indicated by the temperature of 5% weight loss in TGA thermograms. The temperature of 10% weight loss was in the range of 475-490[degrees]C, and char yield at 700[degrees]C was of 23.0-29.3%. The degradation process exhibited a maximum of decomposition ([T.sub.max]) in the interval of 530-540[degrees]C. It can be noted that the polymers 4a and 4b, containing phthaloni-trile groups, exhibited slightly lower initial decomposition temperature when compared with polymer 4c, not containing these groups. Also, the char yield at 700[degrees]C increased with the increase of phthalonitrile content. Thus, polymer 4a, derived from diamine 1, exhibited the highest char yield at 700[degrees]C (29.3%), whereas the polymer 4c without phthalonitrile groups exhibited lower char yield (23.0%).

TABLE 2. Thermogravimcinc parameters for polymers 4,
in air atmosphere.

Polymer       [T.sub.       [T.sub.       [T.sub.  Char yield
             5.sup.a]     10.sup.b]    max.sup.c]      at 700
         ([degrees]C)  ([degrees]C)  ([degrees]C)  (degrees)C

4a                427           475           540        29.3

4b                430           480           530        25.1

4c                450           490           535        23.0

(a.) Initial decomposition temperature, i.e., the temperature
of 5% weight loss.

(b.) Temperature of 10% weight loss.

(c.) Maximum polymer decomposition temperature.

The effect of polar phthalonitrile group concentration on electrical properties was evaluated on the basis of relative permittivity and dielectric loss and their variation with frequency and temperature. In an alternating electric field, the permittivity is a complex quantity, [epsilon]* = [epsilon]' - i[epsilon]", where [epsilon]' is the relative permittivity or dielectric constant, and [epsilon]" is an imaginary component, called dielectric loss or dissipation factor. The three-dimensional diagrams of the frequency and temperature dependency of the [epsilon]' and [epsilon]" for the polymer film 4b, in the first and second scan, are shown in Fig. 4. In the first scan, [epsilon]' increased up to 40[degrees]C; a decrease of [epsilon]' appeared in the temperature range of 40-120[degrees]C, and thereafter [epsilon]' increased with the increase of the temperature. The variation of e' values is related to the net dipolar moment align on the electric field director; therefore, the decrease of El with increasing temperature, in the interval of 40-120C, could be due to the elimination of polar molecules, such as water. In the second scan, [epsilon]' increased continuously with the increase of the temperature over the entire interval of temperature used for measurements. The dependence of [epsilon]" versus temperature and frequency revealed, in the first scan, a [gamma] relaxation process. Such relaxation was observed for all the samples. In the case of the dried samples (second scan), the [gamma] relaxation process was not observed. This suggested the importance for [gamma] relaxation process of the water molecule presence in the polymer films. At higher temperature range, a weak [beta] relaxation can be observed only in the low frequency interval. The values of [epsilon]' and [epsilon]" sharply increased especially at high temperature and low frequency, because of conductivity process. It can be observed that by increasing the frequency, the maxima of the transitions [gamma] and [beta] and of the conductivity process shifted to higher temperatures.

The [epsilon]' values of the polymer films 4, at 10 kHz and various temperatures, in the second scan of the measurements, were in the range of 3.01-3.43 (Table 3). Polymer 4c, not containing phthalonitrile groups, exhibited lower [epsilon]'. By increasing the concentration of these groups, in the case of 4a and 4b. an increase of the [epsilon]' appeared, because of the polar nature of the nitrile groups present in these polymers. Additionally, as it was shown earlier, 4a and 4b exhibited higher water absorption, which also led to higher [epsilon]' values. It was previously reported that [epsilon]' is a function of the total polarizability of a polymer. In spite of a large number of polar amide, imide, and phthalonitrile groups, the polymers 4 exhibited relatively low [epsilon]'. The [epsilon]' values were comparable with that of H polyimide film--a polyimide based on pyromellitic dianhydride and 4,4'-diaminodiphenylether, which is one of the most common high-performance dielectrics used in microelectronic applications, having a dielectric constant of 3.5 (1). This behavior can be explained by the low polarizability of fluorine atoms. The fluorine substitution increases the hydrophobicity thus reducing the absorption of moisture and consequently decreasing [epsilon]'. The 6F groups can also improve the dielectric performance because of less efficient chain packing and increased free volume of polymer (31), (32).

TABLE 3. The dielectric permittivity at 20[degrees]C and 10
kHz, and the activation energy for [gamma], [beta],
and [sigma] transitions of polymers 4.

Ploymer      Dielectric  [E.sub.a] of   [E.sub.a] of   [E.sub.a] of
           permittivity       [gamma]         [beta]        [sigma]
                           relaxation     relaxation     relaxation
                         (first scan)  (second scan)  (second scan)
                            (kJ [mol.      (kJ [mol.      (kJ [mol.
                             sup.-1])       sup.-1])       sup.-1])

4a                3.43          50.5             164            279

4b                3.19          45.9             171            239

4c                3.01          45.1             116            210

The activation energy for [gamma] and [beta] relaxation processes was calculated by applying the Arrhenius equation:

f = A exp(-[E.sub.a]/RT), (4)

where f is the frequency, A is the pre-exponential factor, Ea is the activation energy, R is the gas constant, and T is the peak-maximum temperature. [E.sub.a] of [gamma] relaxation, characteristic on first scan, calculated with Arrhenius equation, exhibited similar values, being in the interval of 45.150.5 kJ [mol.sup.-1], suggesting the occurrence of the same process for this transition. This relaxation was dominated by the presence of small molecules of water and was observed only in the first scan of measurements. The [beta] transition, observed clearly in the second scan of measurements, is generally associated with local bond rotations along the polymer backbone. In general, these motions are considered to be primarily a function of the polymer structure, and their presence and magnitude have been ascribed to several material properties. The [beta] relaxation also demonstrated an Arrhenius behavior. The activation energy of the [beta] relaxation was in the range of 116-171 kJ [mol.sup.-1] (Table 3). The values of activation energy for polymers containing phthalonitrile groups were higher than that of the polymer 4c not containing these groups.

Starting around 150-200[degrees]C, there was a sharp increase in [epsilon]' and [epsilon]". This large frequency-dependent contribution to the dielectric response, especially at low frequencies, may come from an accumulation of charges at the interfaces between regions of different electrical and dielectric properties. Because of the motion of charge carriers, there will be a decay of the applied electric field termed electric field relaxation or conductivity relaxation (33), noted [sigma]. Considering the charges as the independent variable, conductivity relaxation effects can be suitably analyzed within the dielectric modulus formalism in terms of a dimensionless quantity, M*(1). Figure 5a and b presents the dependence of M' and M" versus temperature at 0.l Hz for the sample 4b, respectively. The dispersions of M' and M" indicated a presence of the distribution of conduction relaxation times. In M" diagram, at low field of frequency and high temperature, a peak corresponding to the conductivity process appeared. This peak also shifted to higher temperature with the increase of the frequency. It can be regarded as related to the mobility of charge carriers and leading to system conductivity. This conductivity process masks the primary relaxation connected with glass transition. The activation energy calculated with Arrhenius equation was in the range of 210-279 kJ [mol.sup.-1] (Table 3).

The polymer films 4 were analyzed to obtain nanometric displacements, when an electric voltage is applied on their surface. The measurements were performed with flexible thin electrodes using an interferometric AGI-LENT 5529A system, an instrument dedicated to investigate and measuring the linear microdisplacement and nanodisplacement with a remarkable resolution of 2 nm (34). A plane thin electrode can ensure a good electric contact with the polymer film and good mechanical roughness at the contact with the retroreflector of the equipment. The force exerted on the film was 24 cN and summarizes the weights of the retroreflector and the thin electrode. The specific actuation parameters at 220 V are collected in Table 4: the electric field, E (V [m.sup.-1]) = U/g, where g is the thickness of the film and U is the dc voltage; the maximum value of nanoactuation or nanodis-placement. [d.sub.max](nm): the sensitivity of nanoactuation, S (nm [V.sup.-1]) = [d.sub.max]/U. An improvement of nanoactuation was observed by increasing the phthalonitrile group content. Thus, the sample 4a having the highest content of phthalonitrile groups showed the highest displacement value, when comparing with sample 4c that does not contain such groups.

TABLE 4. Actuation parameters of polymer films 4.

Sample      g (a)      E (b) (V    [d.sub.max] (c)      S (d) (nm
             (mm)  [mm.sup.-1])               (nm)    [V.sup.-1])

4a           0.45          4888               -450           2.04

4b           0.04          5500               -300           1.36

4c           0.06          3666               -120           0.55

(a.) The thickness of the film.

(b.) The electric field.

(c.) The maximum value of nanoactuation or nanodisplacement.

(d.) The sensitivity of nanoactuation.

Figure 6 presents the nanoactuation results for sample 4a, as an example. A manual switch was used to control the dc electric field. The signals that are visualized represent the film actuation response to an electric dc field at different voltage. It is considered that the nanoactuation has an electrostriction character. Electrostriction is a polarization process exhibited by all dielectric materials when an electric field is applied to a dielectric material. This is usually accompanied by a deformation of the material where the strains are proportional to the square of polarization. The types of polarization causing the electro-strictive effect are electronic, ionic, and dipolar polarization. Dipolar polarization is due to reorientation of dipoles under an applied field. The nanodisplacement field increased with the increasing of the electric voltage. Thus, at 130, 180, and 220 V dc, the maxima nanometric displacements were of 150, 300, and 450 nm, with the corresponding sensibility of 1.15, 1.66, and 2.04 nm [V.sup.-1], respectively. A very good actuation behavior was obtained at 220 V. Figure 6c shows four experiments made with different time exposures to the electric field of inverse polarization. The experiments evidenced a time displacement of the reference level, when the voltage is 0, which corresponds to the beginning of the experiment (the reference level is superposed on the axe 0.0 of the AGILENT interferometric measuring system) of approximately--110 nm at 45.5 s.

Figure 7 shows the dependence of nanodisplacements versus load force, to a constant voltage of 220 V dc, for samples 4a and 4c. The load force is represented by the force that acts on the surface of film in the time of electrostrictive experiment. It can be noted that for a load force in the range of 3-7 N. the sample 4a containing phthalonitrile groups exhibited better results than sample 4c not containing these groups. These results are very important for potential application of the present polymers in actuation elements fabrication.


The incorporation of pendant phthalonitrile groups in the macromolecular chains of poly(amide imide)s leads to polymers having good solubility in organic solvents while retaining high thermal stability. The polymers could be processed from their solution into thin flexible films. They were thermally stable up to 420[degrees]C. The presence of polar phthalonitrile groups increased the water absorption and dielectric constant. The polymer films showed nanoactuation in the range of 120-450 nm at 220 V dc, with flexible thin electrodes, depending on the phthalonitrile group content.


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Elena Hamciuc, (1) Corneliu Hamciuc, (1) lonela-Daniela Carja, (1) Tachita Vlad-Bubulac, (1) Mircea Ignat (2)

(1.) "Petru Pon!" Institute of Macromolecular Chemistry, Aleea Gr. Ghica Voda 41A, Iasi 700487, Romania

(2.) National Institute for Research and Development in Electrical Engineering ICPE-CA, Splaiul Unirii 313, Bucharest 030138, Romania

Correspondence to: Elena Hamciuc; e-mail: Contract grant sponsor: CNCSIS--UEFISCDI; contract grant number: PNII--IDE1 code ID_997/2008.

Published online in Wiley Online Library (

[c] 2012 Society of Plastics Engineers

DOI 10.1002/pen.23268
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Author:Hamciuc, Elena; Hamciuc, Corneliu; Carja, lonela-Daniela; Vlad-Bubulac, Tachita; lgnat Mircea
Publication:Polymer Engineering and Science
Article Type:Report
Geographic Code:4EXRO
Date:Feb 1, 2013
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