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Molecular Dynamics in Nanostructured Polyimide–Silica Hybrid Materials and Their Thermal Stability V. A. BERSHTEIN,1 L. M. EGOROVA,1 P. N. YAKUSHEV,1 P. PISSIS,2 P. SYSEL,3 L. BROZOVA4 1

Ioffe Physico-Technical Institute, Russian Academy of Sciences, 194021 St. Petersburg, Russia

2

Department of Physics, National Technical University of Athens, Zografou Campus, 15773 Athens, Greece

3

Institute of Chemical Technology, Technicka 5, 16628 Prague 6, Czech Republic

4

Institute of Macromolecular Chemistry, Heyrovsky square 2, 16206 Prague, Czech Republic

Received 4 April 2001; revised 21 February 2002; accepted 21 February 2002 Published online 00 Month 2002 in Wiley InterScience (www.interscience.wiley.com). DOI: 10.1002/polb.10162

Molecular motion and thermal stability in two series of nanophase-separated polyimide–silica (PI–SiO2) hybrid materials with chemically bound components were studied. The hybrids were synthesized from p-aminophenyltrimethoxysilane-terminated poly(amic acid)s as PI precursors and tetramethoxysilane as a silica precursor via a sol– gel process. The hybrids differed in their PI chemical structure and chain length (numberaverage molecular weight ⫽ 5.000, 7.500, or 10.000) and in their SiO2 content, which ranged from 0 to 50 wt %. Differential scanning calorimetry, laser-interferometric creep rate spectroscopy, and thermally stimulated depolarization current techniques were used for studying the dynamics from 100 to 650 K and from 10⫺3 to 10⫺2 Hz. Comparative thermogravimetric measurements were also carried out from 300 to 900 K. Silica nano- or submicrodomains that formed affected PI dynamics in two opposite directions. Because of the loosening of the molecular packing of PI chains confined to nanometer-scale spaces between silica constraints, an enhancement of small-scale motion, mostly at temperatures below the ␤-relaxation region, occurred. However, a partial or total suppression of segmental motion could be observed above the ␤-relaxation temperature, drastically so for the shortest PI chains at elevated silica contents and within or close to the glass-transition range, because of the covalent anchoring of chain ends to silica domains. Large changes in thermal stability, including a 2.5-fold increase in the apparent activation energy of degradation, were observed in the hybrids studied. A greater than 100 °C rise in long-term thermal stability could be predicted for some hybrids with respect to pure PI. © 2002 Wiley ABSTRACT:

Periodicals, Inc. J Polym Sci Part B: Polym Phys 40: 1056 –1069, 2002

Keywords: nanocomposites; polyimides; molecular dynamics; glass transition; thermal properties

INTRODUCTION In recent years, steadily increasing attention has been paid to organic–inorganic hybrid materials

Correspondence to: V. A. Bershtein (E-mail: vbersht@ polmater.ioffe.rssi.ru) Journal of Polymer Science: Part B: Polymer Physics, Vol. 40, 1056 –1069 (2002) © 2002 Wiley Periodicals, Inc.

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prepared via the sol– gel process, especially polymer–silica hybrids obtained by the hydrolysis and condensation of tetraalkoxysilane, tetramethoxysilane (TMOS), or tetraethoxysilane (TEOS) in the presence of a polymer or oligomer or some polymer precursor.1–20 Of special interest are highly homogeneous, optically transparent hybrids with nanometer-scale silica domains (clusters) covalently bound with a polymer. They may

POLYIMIDE–SILICA HYBRID MATERIALS

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be prepared by the incorporation of trialkoxysilyl moieties into polymers either as the end (endcapping agents) or as pendant groups. When polycondensation between TEOS (or TMOS) and trialkoxysilyl groups takes place, chemical bonds between polymer and silica nanophases are formed. This procedure was recently used for preparing different polymer–silica hybrid materials with 10 –100-nm silica domains, such as from trialkoxysilane-terminated polydimethylsiloxane,1,3 poly(arylene ether) ketone,4 poly(phenylene terephthalamide),5 poly(styrene-co-maleic anhydride),6 ethylene–propylene– diene terpolymer,7 poly(⑀-caprolactone), and 8 polyimides (PIs).9 –13 Moreover, polymers deprived of groups reactive in the sol– gel process but prone to hydrogen bonding with silanol groups have also been successfully incorporated into nanophase-separated hybrid materials with a high degree of mixing of their components.8,14 –18 Modifying known polymers with inorganic nanoclusters promises improvements in some polymer properties through the possible tailoring of the solid-state properties, depending on the nature and relative content of the constitutive polymer and inorganic components. Therefore, great potential is expected for hybrid materials in microelectronics9 and membrane technology21,22 and as optical wave guides,23,24 biosensors,25 and biocompatible8 and aerospace9 materials. The hybrids may be considered some of the next generations of polymeric materials for advanced technologies. Among the hybrid materials, nanostructured PI–silica materials are especially interesting because of the commercial importance of PIs due to their chemical resistance and good mechanical and dielectric properties at elevated tempera-

tures. As is well known, the physical properties and free-radical reactions of thermooxidative degradation in solid polymers are controlled to a large extent by segmental motion. Meanwhile, information on the dynamic behavior of organic– inorganic hybrid materials is very scarce. Only a slight increase in the glass-transition temperature (Tg) was observed for flexible-chain ethylene–propylene– diene terpolymer,7 poly(⑀-caprolactone),8 and poly(amide-imide)15 in hybrids. In our preliminary communication on PI–silica hybrids,26 the essential changes in the segmental dynamics and an improvement in the thermal stability of high-performance PIs were shown. The purpose of this article is to present in detail the results of experimental research on PI– silica hybrid materials that was aimed at answering two questions: (1) how do the silica nanoclusters, introduced via a sol– gel process, influence PI molecular dynamics measured over broad temperature ranges and (2) what changes in thermal stability may occur in the PI–silica hybrids with respect to pure PI?

were obtained as PI1 and PI2 precursors, respectively. In turn, the most reactive tetraalkoxysilane, TMOS, was used as the starting compound for the noncatalyzed sol–gel process, that is, as a silica precursor.

During the second stage, a calculated amount of TMOS was added to a 10 wt % solution of the aforementioned end-capped poly(amic acid) in N-methyl-2-pyrrolidone, and then a stoichiomet-

EXPERIMENTAL Materials Two series of PI–silica hybrids, differing in their PI molecular structures, were synthesized via a sol– gel process as earlier described.13,26 During the first stage, the starting compounds, p-aminophenyltrimethoxysilane-terminated poly(amic acid)s based on 4,4⬘-diaminodiphenylether and 3,3⬘oxydiphthalic anhydride (I) or 4,4⬘-(hexafluoroisopropylidene)diphthalic anhy‘dride (II),

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ric amount of water per methoxy group of TMOS was added that led to the hydrolysis of all methoxysilane groups. The resulting homogeneous solution was poured onto a Teflon dish. After drying at 60 °C for 24 h, the film was heated 2 h at 125 °C, 2 h at 200 °C and 5 h at 300 °C. During heating, both silanol condensation and poly(amic acid) imidization reactions proceeded, resulting in PI–silica hybrid formation in which silica domains were covalently bound with the ends of PI chains. The silica content in hybrids was calculated under the assumption that the following reactions Si共OCH 3兲 4 ⫹ 4H 2O 3 Si共OH兲 4 ⫹ 4CH 3OH Si共OH兲 4 3 SiO 2 ⫹ 2H 2O proceeded to completion. However, combined density, thermogravimetric analysis (TGA), and IR spectroscopy measurements showed that the aforementioned thermal treatment (the final stage was limited to 5 h at 300 °C to avoid PI thermal degradation) was insufficient to complete the sol– gel reaction and form perfect silica clusters. As a result, the samples obtained contained, according to IR data, about 5 wt % potential volatiles.13 This means that silica domains contained a small quantity of unreacted hydroxyl groups; that is, the transformation of the silica precursor to silica did not reach 100% under these conditions. Such silica networks are characterized by sharply reduced Tg values, which are equal to about 500 –700 °C instead of 1200 °C for typical amorphous silica.27 At the same time, IR spectroscopy measurements showed that silica phases did not prevent the imidization process in hybrids, which was basically completed after a long treatment at 300 °C.12 Two series of transparent hybrid films, PI1– silica and PI2–silica, 0.1– 0.3 mm thick with various compositions were obtained. On the whole, 23 samples with regularly varied compositions and the common formulas PI1–m/n and PI2–m/n were studied (Table 1). Using such a large number of samples allowed us to follow the changes in the molecular dynamics of these hybrid materials with four factors: (1) PI chain rigidity, (2) PI chain length, (3) silica content, and (4) the presence or absence of a hybridization effect. Unlike PI1 chains containing many oxygen atom hinges, PI2 chains are more rigid because of the replacement of the majority of oxygen atoms

Table 1. Polyimide–Silica Hybrid Films Sample PI1–5/0 PI1–5/10 PI1–5/20 PI1–5/30 PI1–5/40 PI1–5/50 PI1–7.5/10 PI1–7.5/20 PI1–7.5/30 PI1–7.5/50 PI1–10/10 PI1–10/30 PI1–10/40 PI1–10/50 PI1–5/10 blend PI2–7.5/0 PI2–7.5/10 PI2–7.5/12.5 PI2–7.5/17.5 PI2–7.5/20 PI2–7.5/24 PI2–7.5/30 PI2–7.5/10 blend

Mn (g mol⫺1) Silica Content (wt %) 5000 5000 5000 5000 5000 5000 7.500 7.500 7.500 7.500 10.000 10.000 10.000 10.000 5.000 7.500 7.500 7.500 7.500 7.500 7.500 7.500 7.500

0 10 20 30 40 50 10 20 30 50 10 30 40 50 10 0 10 12.5 17.5 20 24 30 10

with C(CF3)2 groups. In the aforementioned formulas, m is equal to 5, 7.5, or 10, corresponding to number-average molecular weights (Mn) of 5.000, 7.500, or 10.000 g mol⫺1, respectively, for PI chains, whereas n is a silica content in the samples that ranged from 0 to 50 wt %; that is, pure PIs were also studied for comparison. For the estimation of the influence of the hybridization effect itself, PI–silica blends, prepared from TMOS and non-end-capped materials, usually poly(amic acid) (i.e., without the introduction of functional end groups and the formation of chemical bonds between both components), were also studied. In previous experiments,12,13 scanning electron micrographs were obtained for some of these samples on cross sections of the films covered with a thin gold layer with a JEOL 100B electron microscope. In Figure 1, two examples of these micrographs are given for the PI1–silica hybrid and a blend with an identical ratio of both components. The blend was a microphase-separated product with silica particles of about 1␮m [Fig. 1(a)], whereas the hybrid film was nanostructured with silica clusters of about 100 nm [Fig. 1(b)]. As shown,12,13 some changes in the structures of PI– silica hybrids took place, depending on the silica


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Figure 1. Scanning electron micrographs of (a) the PI1–5/10 blend and (b) the PI1– 5/10 hybrid material and (c) a simplified scheme of the hybrid structure. The bar is equal to 1 ␮m.

content. Nevertheless, the common tendency of the formation of nanostructured hybrids and microstructured blends was maintained. Figure 1(c) shows a simplified scheme of the PI–silica hybrid structure in which PI chain ends are attached covalently to silica nanoclusters. Characterization Four complementary techniques— differential scanning calorimetry (DSC), laser-interferometric creep rate spectroscopy (CRS), thermally stimulated depolarization current (TSDC) techniques, and TGA—were used for the characterization of

molecular dynamics and thermal stability in the hybrids studied. A DSC-2 PerkinElmer apparatus was used for the characterization of the glass transition, namely, Tg, at the half-height of the heat capacity step; the transition range (⌬Tg ⫽ Tg⬙ ⫺ Tg⬘, where Tg⬘ and Tg⬙ are the onset and completion temperatures of the glass transition); and the values of the heat capacity step (⌬Cp) and the activation energy (Q) of segmental motion within the transition range. The experiments were carried out at a heating rate (v) of 5, 10, 20, or 40 K min⫺1 in a nitrogen atmosphere from 300 to 640 K.

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The reproducible DSC curves, obtained during the second and following scans, that is, after water removal and sample stabilization, with the samples cooling from 640 to 300 K at v ⫽ 320 K min⫺1, were used for the estimations. The Q values were determined from the displacement of Tg, Tg⬘, or Tg⬙ with the heating rate according to the following formula:28 Q ⫽ ⫺ Rdlnv/d共T ⫺1兲

(1)

For estimation over a wide range of frequencies, the activation energy of the cooperative glass transition is generally variable because of the dependence of segmental motion cooperativity on frequency, especially at high frequencies. However, under our experimental conditions, at the correlative (equivalent) frequency (␯corr ⬇ 10⫺2 Hz), that was not the case after approximation because plots of ln v versus Tg⫺1 turned out to be linear. The accuracy of Tg, ⌬Cp, and Q estimates were equal to about 1 K, 3%, and 10%, respectively. CRS and TSDC techniques were used for an analysis of molecular motion in the hybrids below T g. The laser-interferometric CRS technique has been described elsewhere.29 –31 Here, the creep rate versus the temperature dependence (creep rate spectrum) was obtained from 290 to 520 K at a low tensile stress (␴ ⫽ 1 MPa), that is, on the basis of a deformation increment of only about 0.01% (a few hundred nanometers). Besides the stress, the second experimental parameter, the time (t ⫽ 10 s) from the moment of loading to the onset of measuring, was kept constant at different temperatures. The instrumental error of creep rate measurements did not exceed 1%. The correlative frequency of the CRS experiments, ␯corr ⬇ 10⫺2 Hz, was virtually the same as in the DSC measurements. The TSDC method, which measures the thermally activated release of stored dielectric polarization, corresponds to measuring dielectric losses against temperature at low equivalent frequencies of 10⫺4 to 10⫺2 Hz.32 The TSDC apparatus used has been described in detail elsewhere.33 Samples in the form of 14-mm-diameter discs, cut from films 0.1– 0.3 mm thick, were used. Each sample was inserted between brass electrodes (capacitor plates) and polarized by the application of an electric field (Vp ⫽ 100 V) for the time tp ⫽ 5 min at the temperature Tp ⫽ 25 °C. Then, the

Figure 2. Creep rate spectra obtained for PI1–5/0 and PI1–5/20 hybrid materials under extension and indicated ␴ and t values.

sample was cooled under the electric field at a rate of 6 K min⫺1 down to 100 K. During the short-circuiting of this sample and heating at a rate of 3 K min⫺1, a plot of the discharge current versus temperature (TSDC spectrum) was measured from 100 K to room temperature. TGA measurements were carried out with a DuPont 990 thermal analyzer (module 951) during heating of the PI1 and PI1–silica hybrid films in air from 20 to 640 °C at rates of 5, 10, or 20 K min⫺1.

RESULTS AND DISCUSSION Molecular Dynamics below Tg Typically, broad ␤- and ␥-relaxation regions are observed in PI, for which the ␤ relaxation is associated with segmental motion in loosely packed PI domains28,34,35 and the ␥ relaxation has been attributed to noncooperative motion of imide cycles35 or to the influence of sorbed water.36 Figure 2 shows the creep rate spectra obtained from room temperature to 520 K for pure PI1 and PI1–silica hybrid materials with the shortest chains (Mn ⫽ 5.000). Nanometer-scale silica domains covalently bound with PI chain ends may influence molecular dynamics below Tg in two opposite directions: (1) decreases in the creep rates at temperatures above the ␤-relaxation temperature (T␤ ⬇ 350 –370 K), including the region of so-called intermediate relaxations with the maximums at about 420, 440, and 460 K and


particularly drastically at 480 –520 K, in the vicinity of Tg, and (2) increases in the creep rates measured at T ⬍ T␤. Small displacements of the intermediate peaks at 440 and 460 K to higher temperatures are also observed. According to the experimentally proven concept of common segmental origin of ␣ and ␤ relaxations in flexible-chain polymers,28,37–39 these relaxation transitions are associated with intermolecularly cooperative or noncooperative rotary movement of Kuhn segments, respectively. Therefore, Figure 2 allows us to assume that silica nanodomains in the hybrids may suppress, to some extent, just segmental motion (especially cooperative motion) but facilitate small-scale movement. As shown later, the DSC and TSDC experimental data obtained are consistent with this assumption, and these results turned out to be readily interpreted on the basis of this concept and nanoconfinement and constraining (chainanchoring) effects. Figure 3 and Table 2 show the results of TSDC measurements carried out from 100 to 290 K for PI2 and its hybrids with 10 and 20% silica under three different humidity conditions (levels of water uptake). The dry samples were prepared by being kept for 52 h in a vacuum of 5 ⫻ 10⫺3 Torr at 150 °C. In addition, moderately moistened samples (under ambient conditions) and samples water-saturated in an atmosphere of 96% relative humidity were studied. On the whole, an overall enhancement of small-scale molecular mobility was found due to the incorporation of the silica nanodomains into PI. The other result is that the hybrids absorbed more water than pure PI, and water uptake substantially increased with silica content in a hybrid (Table 2). Three TSDC peaks are observed below room temperature, including (1) the ␥-relaxation peak at about 130 K; (2) the peak at 260 –280 K corresponding, obviously, to a sub-T␤ relaxation at 300 K in the creep rate spectrum of PI1 (Fig. 2); and (3) a narrower interfacial peak at about 190 K observed for the hybrid with 20% silica (Fig. 3). As seen in Table 2, there is no appreciable temperature shift of the ␥-relaxation TSDC peak with sample composition, but the ␥ and interfacial peaks shift slightly, by 3–5°, to lower temperatures in the water-saturated samples. At the same time, the peak magnitude increases in the hybrids to a great extent with silica content. Therefore, about a threefold increase in the ␥-relaxation peak is observed at all levels of water

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Figure 3. TSDC thermograms measured for PI2– 7.5/0, PI2–7.5/10, and PI2–7.5/20 films in the temperature region below the ␤-relaxation range under different humidity conditions: (a) in an atmosphere of 96% relative humidity, (b) under ambient conditions, and (c) in the dry state. See also the Table 2.

uptake in the samples, including the dry ones. The latter indicates that this relaxation is not associated indispensably with the presence of sorbed water in a sample. However, water affects to some extent the relaxation strength of the TSDC peaks under study. The most pronounced influence of water takes place for a sub-T␤ relaxation peak when sorbed water is in a liquid, mobile state. Water penetration inside silica domains must be negligibly small. Therefore, the effect of increased water uptake in the hybrids suggests enhanced water accumulation in polymeric domains and, maybe, partly at the internal interfaces between both components within interfacial polymer layers and at the silica surface. Such increased hybrid plasticization by water may be considered a proof of loosened molecular packing and excess free volume in the hybrids with respect to pure PI. It may be assumed that such loosening originated from some steric hindrances to dense mo-

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Table 2. Characteristics of ␥-Relaxation and Interface Peaks in TSDC Spectra of PI and Hybrids under Different Wetness Conditions

Dry Conditions

96% Relative-Humidity Conditions

Room Conditions

Sample

Tm (K)

Im (pA)

Tm (K)

Im (pA)

W%a

Tm (K)

Im (pA)

W%a

PI2–7.5/0 PI2–7.5/10 PI2–7.5/20

134 134 133 193b

0.31 0.65 1.10 0.7b

132 134 132 194b

0.35 0.61 1.25 0.9b

0.84 1.67 1.97

129 130 128 190b

0.65 0.88 1.75 1.05b

1.68 4.17 4.92

a b

Water uptake Interface.

lecular packing for two reasons: (1) the covalent anchoring of the chain ends to silica domains and (2) the polymer location in nanoscale spaces between these domains (rigid walls). It is associated with the common, well-known problems of nanoscale confinement and substance– constraint interactions in molecular dynamics, and we return to this problem later when we discuss glasstransition behavior in these hybrids. The new peak, manifesting itself in the TSDC spectra of the hybrid with sufficiently high (20%) silica content only, cannot be attributed to Maxwell–Wagner–Sillars interfacial polarization, often observed in heterogeneous systems, typically at high temperatures and low frequencies.32,33 The reason is that the peak occurs here at low temperatures, at which the components do not show any significant conductivity. The peak is tentatively attributed to small-scale dynamics of the ␥-relaxation type within immobilized segments in the interfacial PI layers (discussed later). An increase in potential barriers to torsional vibrations of imide cycles may just provide a shift of the ␥-relaxation peak from 130 to 190 K in the TSDC spectra (Fig. 3).

blend with silica microparticles [see Fig. 1(a)], and the nanostructured hybrids differing in PI1 molar mass and silica content. Sample stabilization was attained by preliminary heating of the samples to 640 K. In such cases, the second and subsequent scans showed identical DSC curves. Figure 5 shows the compositional dependence of glass-transition characteristics, including ⌬Tg, Tg, and ⌬Cp values. The data given in Figures 4 and 5 are as follows. For the PI1–5/10 blend, when there was no chemical bonding between the components, Tg slightly decreased but ⌬Tg broadened from 7 ⫾ 1°C in pure PI1 to 20°C in this blend, basically toward lower temperatures. For PI1–silica hy-

Glass Transition Figures 4 –7 represent the results of a DSC study of the glass transition in both series of PI–silica hybrids listed in Table 1. Rather drastic changes in the characteristics of the PI glass transition can be observed after the incorporation of the silica nanodomains chemically bound with a polymer. Figure 4 shows some typical DSC curves obtained over the temperature region of the glass transition for stabilized samples of pure PI1, its

Figure 4. DSC curves obtained in the glass-transition region for the PI1–5/0, PI1–5/10 blend, and indicated hybrids differing in PI1 chain length and silica content. The second scan was taken after heating to 640 K at a rate of 20 K min⫺1 and subsequent cooling to 300 K at a rate of 320 K min⫺1. The heating rate was 20 K min⫺1.


Figure 5. DSC data for the PI1–silica hybrids with different PI1 molar masses: Tg, ⌬Tg, and ⌬Cp as functions of the silica content. The heating rate was 20 K min⫺1. The dashed line corresponds to a calculated decrease in ⌬Cp on the basis of additivity with the silica content.

brids with the longer chains (Mn ⫽ 10.000 or 7.500), the most pronounced broadening of the transition range occurred, and this effect increased with silica content. Therefore, ⌬Tg rose up to 35°C in breadth for the PI1–7.5/50 sample, and the transition broadened to both higher and lower temperatures. Tg⬘ shifted from 540 to 528 K, whereas Tg⬙ increased from 548 K to 563 K. At the same time, Tg increased weakly, and a somewhat nonadditive decrease in ⌬Cp values in the hybrids was observed [Fig. 5(c)]. The glass transition in the hybrids with the shortest PI1 chains (Mn ⫽ 5.000) behaves in a different way (Figs. 4 and 5). In this case, Tg increased from 544 to 565 K, and the broad glasstransition range was entirely displaced to higher temperatures; Tg⬙ increased up to 580 K. In addition, a pronounced, anomalous decrease in ⌬Cp with the silica content was observed. At 40 –50% silica in these hybrids, no heat capacity step was

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observed in the DSC curve over the temperature range under study (up to 640 K until the onset of polymer degradation). This indicates total suppression of cooperative segmental motion in the glass transition. Consequently, the effect of segmental motion suppression turned out to be maximal for the hybrids with the shortest PI1 chains. Figure 6 shows the similar compositional dependencies of the glass-transition parameters determined for the second series of the hybrids, with Mn ⫽ 7.500 for PI2 chains with increased rigidity. Qualitatively similar results, as in Figure 5, were obtained here. However, the drop in the ⌬Cp step is larger for the PI2-containing hybrids than for the PI1-containing ones with the same silica content and chain length. This means that the effect of segmental motion suppression manifests itself to a larger extent in the hybrids with more rigid macromolecules. Calorimetric measurements, carried out through heating at different rates for the series of identical, totally stabilized samples, allowed us to estimate the magnitudes of the effective activation energy (Q) for segmental motion within the glass transition as calculated from the displacement of Tg, Tg⬘, or Tg⬙ with the heating rate. Some examples of experimental plots of ln v versus T⫺1 and Q values calculated therefrom are shown in Figure 7. A high level of Q values, typical of cooperative motion in the PI glass transition,28 was

Figure 6. DSC data for PI2–silica hybrids: Tg, ⌬Tg, and ⌬Cp as functions of the silica content. The heating rate was 20 K min⫺1. The dashed line corresponds to a calculated decrease in ⌬Cp with the silica content, whereas the arrow indicates blend behavior.

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Figure 7. DSC data: ln v versus the reciprocal of Tg and Q calculated therefrom for (1) PI2–7.5/0 (for Tg), (2) PI1–5/30 (for Tg⬙), (3) PI1–5/30 (for Tg⬘), (4) PI1–5/10 (for Tg⬘), and (5) PI1–5/0 (for Tg⬙).

also retained in the hybrids. However, both reduced (from 600 – 620 to 520 kJ mol⫺1) and increased (from 720 to 840 kJ mol⫺1) activation energies were obtained. Moreover, the latter values were obtained for two temperatures inside the glass-transition range (in the PI1–5/30 sample; see Fig. 7), which indicates some dynamic heterogeneity near Tg in such a case. The effects described, characterizing the anomalous glass-transition behavior in PI–silica hybrid materials (nanocomposites), may be interpreted in terms of (1) the special dynamic behavior of a polymer confined to nanoscale or submicroscopic spaces caused by the changes in molecular packing (pure confinement effect) and (2) the influence of strong polymer–rigid constraint interactions (the covalent anchoring of chain ends to silica nanodomains, in our case). Both influences relate to the topical confinement– constraining problem in molecular dynamics. Much attention is currently being paid to studying the peculiarities of dynamics around Tg in glass-forming polymers and simple liquids confined to nanoscale spaces and at interfaces. Their special dynamic properties in nanopores or in thin, 10 –50-nm-thick polymer films, deposited on solid substrates or freely standing or located between foreign materials, have been found (see refs. 40 –53). Under these conditions, a sharp broadening of ⌬Tg and Tg displacement to both higher and lower temperatures were observed. Therefore, for polymer films less than or approximately equal in thickness to one unperturbed radius of gyration or a few Kuhn segments, Tg reduction down to T␤ was shown,43 whereas film grafting to or hydrogen bonding with a silica surface led to a Tg shift upward.43,44

Furthermore, when studying these finite-size effects, some heterogeneity in dynamic behavior around Tg was revealed that was associated with some difference in the state of interfacial and more remote domains of confined material. That was interpreted in terms of two-layer47,48 or three-layer models.49,50 Experimental proofs were obtained testifying in favor of the existence of an interfacial layer with partially or totally suppressed mobility (an immobilized or dead layer) in the direct vicinity of rigid constraint as a result of strong substance– constraint interactions. Some estimates indicate that such a layer in substratesupported films50 or in filled compositions53 or including chain sections close to the network junction as a constraint54 may extend to a thickness of 2–5 nm, that is, to one or two Kuhn segments in flexible-chain polymers. Moreover, the inner, more remote layers with either a reduced Tg value or the usual, bulklike behavior are considered in the dynamic models of confined substances.47–50 For instance, there were found, within polystyrene (PS) films 70 –300 nm thick deposited on silicon, an immobilized layer 5 nm thick near the substrate, a 2-nm layer with a reduced Tg value at the polymer–air interface, and an inner layer with bulklike dynamic properties.50 The notion of a confinement effect is quite applicable to an interpretation of the anomalies of segmental dynamics in the majority of modern high-tech polymeric materials: block copolymers, polymer networks, blends, composites, highly crystalline polymers, and organic–inorganic or organic– organic hybrid materials. As a matter of fact, the rigid blocks and rodlike chain constituents or rigid junctions in a network, filler particles, or reinforcing fibers, as well as rigid nanophases and lamellar crystallites, may just play the role of rigid structural constraints. Therefore, a few similar effects of the anomalous glass transition in complex systems were observed, including (1) multifold broadening of the ⌬Tg range; (2) considerable Tg shifts upward or downward to T␤; (3) sharp changes, in both directions, of the activation energy of segmental motion (Q); (4) rising dynamic heterogeneity within the glass-transition range; and (5) threefold to fourfold drops in the motional event scale. These anomalies have been found for highly flexible polydimethylsiloxane (PDMS) blocks in block copolymers with different kinds of rigid blocks,28,55–58 PS chains in fullerene (C60) core PS stars,58 and intercrystalline layers in typical


highly crystalline polymers such as polyethylene (PE).28 The nanoscale confinement effect may best be treated with the concept of the common segmental nature of the glass transition and ␤ relaxation in flexible-chain polymers; these transitions are considered the events of cooperative or local noncooperative motion (at sites of loosened packing) of Kuhn segments, respectively.28,37–39 As found experimentally,28,37,38,59 the degree of intermolecular cooperativity of segmental motion in the glass transition at low frequencies is Z ⫽ 4 ⫾ 1. From this viewpoint, loosened molecular packing of chains and excess free volume, arising in confined nanospaces, result in the partial breakdown or total collapse of intermolecular cooperativity of segmental motion. Therefore, the total collapse of cooperative dynamics with passage to the dynamics of isolated molecules (the transformation of the glass transition into Arrhenius ␤ relaxation, in our interpretation28) was shown recently for a liquid confined to zeolite channels of 0.55 nm.52 The nanocomposites under study are complex polymer systems with the most severe, three-dimensional confinement47,48 of short chains anchored from both ends to rigid constraints. Therefore, all the aforementioned considerations may account for all of the results represented in Figures 4 –7. Indeed, the Kuhn segment length in PIs with oxygen hinges in chains, such as PI1, equals about 5 ⫾ 1 nm.28 This means that macromolecules in the hybrids with the shortest PI1 chains (Mn ⫽ 5.000) consist of three to five segments only, and the majority of their segments must turn out in a more or less immobilized state. This is actually confirmed by the displacement of Tg and the entire ⌬Tg range to higher temperatures, the anomalous large reduction of the ⌬Cp step, and the Q value increasing from about 600 kJ mol⫺1 in pure PIs to 720 and 840 kJ mol⫺1 for the glass transition in the PI1–5/30 hybrid. The total suppression of glass transition in the PI1–5/40 and PI1–5/50 samples (Figs. 4, 5, and 7) shows that the effect of the polymer– constraint interaction predominates in these cases. Therefore, the suppression of cooperative dynamics in the glass transition may be expected for such systems at l/A ⱕ 3–5, where l is the contour chain length and A is the Kuhn segment length. In contrast, the extension of the glass transition to lower temperatures in the PI1–5/10 blend, in the absence of strong polymer– constraint interactions, may be associated with a prevailing

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effect of loosened molecular packing with addition of silica microparticles as fillers. A complicated dynamic behavior took place for the PI1–7.5/50 and PI1–10/50 hybrids with the longer (l/A ⬇ 8 ⫾ 2) polymer chains. The effects of very large extensions of the ⌬Tg range to both higher and lower temperatures with only minor changes in Tg magnitudes and much less suppression of segmental motion (as indicated by the plots ⌬Cp vs composition in Fig. 5) were observed. Such behavior suggests three contributions to segmental dynamics at local sites, including (1) loosened packing (a decreased value of Tg⬘), (2) packing typical of pure PI1 (the usual value of Tg), and (3) somewhat suppressed motion near the point of chain anchoring (an increased value of Tg⬙). Resistance to Thermooxidative Degradation Thermogravimetric measurements were performed for PI with the shortest chains (the PI1– 5/0 sample) and its hybrids with different contents of silica nanodomains. As shown previously, in this series of hybrids, cooperative segmental motion in the glass transition was suppressed to the largest extent, down to the total disappearance of the transition at the DSC curve (Fig. 4). TGA data are given in Table 3. The silica nanophases markedly influenced the decomposition temperatures even in the initial stages of the degradation process at the mass loss ⌬m ⫽ 5– 15%. The magnitude of the degradation effect depended on both the heating rate and the degradation degree of a material. The general tendency here is increasing thermal stability, as estimated by weight loss, with hybridization and with a reduction of the heating rate. At ⌬m ⫽ 5%, the maximum effect was observed for the PI1–5/20 hybrid; however, it is not typical because this initial stage of weight loss depends largely on the evaporation of water and low molecular weight admixtures. As seen from the insert in Figure 8, the onset temperature of chain decomposition, with subsequent evaporation of its fragments, turned out to be, at a heating rate of 5 K min⫺1, about 100 °C higher in the hybrid (T2) than in pure PI1 (T1). In these short-term experiments, degradation temperatures, corresponding to ⌬m ⫽ 10 –15%, increased with silica content in the hybrids, by 50 °C in the limit. A more pronounced positive influence of silica nanoclusters on PI thermal stability was observed during the late stages of the degradation process.

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Table 3. TGA Data Obtained for PI1 and PI1–Silica Hybrid Materials

Weight Loss (%)

Degradation Temperature (°C) at a Heating Rate (K min⫺1) 5

5 10 15

483 520 532

5 10 15

493 530 546

5 10 15

520 539 548

5 10 15

499 572 579

5 10 15

509 576 588

10

20

PI1–5/0 514 548 551 586 563 596 PI1–5/10 516 548 553 578 570 594 PI1–5/20 539 560 558 577 566 586 PI1–5/40 515 530 591 609 596 614 PI1–5/50 523 538 591 607 605 623

Thermooxidative Degradation Activation Energy (kJ mol⫺1)

110 115 115 156 167 169 193 215 215 230 250 250 240 280 255

The data obtained allowed us to estimate, with an accuracy of 10%, the values of the effective activation energy of thermal degradation (Eeff) for the hybrids and pure PI1 with the linear ln v versus T⫺1 dependence and a relation similar to eq 1, where T was the temperature of a definite mass loss of 5, 10, or 15%. Figure 9(a) shows identical slopes of the linear ln v versus T⫺1 plots for degradation over mass losses of 5, 10, and 15%. Therefore, the estimated Eeff values were invariable, within the experimental accuracy of ⫾10%, for each sample at the indicated degradation degrees. At the same time, Eeff increasing with silica content in the hybrids was observed [Fig. 9(b)], and it changed from 110 –115 kJ mol⫺1 for pure PI1 to 240 –280 kJ mol⫺1 for the PI1–5/50 sample (Table 3). Greatly increasing Eeff values in the PI–silica hybrids, as estimated from TGA data, are intriguing experimental facts. However, the interpretation is very difficult. First, thermal decomposition in a solid medium is a complex process associated with different (primary and secondary) chemical reactions, kinetics of reaction product transfer,

and so forth.60 In addition, chemical and diffusion processes in solid polymers are controlled to a considerable extent by segmental dynamics.61 The situation becomes much more complicated in nanocomposites, in which the presence of inorganic phases may influence the kinetics of product transfer and the probability of back and secondary reactions. Therefore, the Eeff values, obtained by TGA, are some empirical characteristics (temperature coefficients) that are only used for comparative estimation of thermal stability. These cannot be considered real physical barriers. In other words, the effect of Eeff increasing may be associated not only with the increasing fundamental thermal stability of PI chains in the presence of silica nanophases but also with the aforementioned accessory factors. Nevertheless, high thermogravimetric values of Eeff are usually associated with thermally stable materials, whereas low values relate to nonstable ones.60 Therefore, three points are of interest in these TGA data: (1) the greatly increasing Eeff values in the hybrids, (2) the regular relation of this effect to silica content in the hybrids (Table 3), and (3)

Figure 8. Thermooxidative degradation of PI1 and PI1–silica hybrids: TGA curves obtained at a heating rate of 5 K min⫺1. The inset shows the weight loss normalized to the sample weight at 200 °C as a function of temperature. The temperatures T1 and T2 correspond to the onset of real chain destruction of the PI1–5/0 and PI1–5/30 samples, respectively.


Figure 9. TGA measurements: ln v versus the reciprocal of T for indicated weight losses in pure PI1 and its hybrids (a) at different loss magnitudes and (b) for different sample compositions.

the connection between Eeff and the degree of suppression of cooperative segmental motion in the glass transition by silica nanodomains (see Fig. 5, a plot of ⌬Cp vs composition). The sharp Eeff rise allowed us to suppose that the most pronounced improvement of the resistance to thermooxidative degradation (with mass loss) in the hybrids could be expected from a very long influence of high temperatures (long-term thermal stability). To predict, tentatively, the magnitude of such a probable effect, we used an approach presented elsewhere.60,62 Generally, the kinetics of thermal degradation may be described by the following formula: dc/dt ⫽ K共1 ⫺ C兲 n

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mass; K ⫽ Z exp(⫺Eeff/RT) is the apparent reaction rate constant; Z is the pre-exponential factor; n is the reaction order; and T is the temperature at which a given mass loss is attained for time t. Assuming, for simplicity, that the degradation process is a first-order reaction and proceeds from the Eeff values, one could calculate Z values and then determine the temperatures at which a given degree of thermal degradation (mass loss) is attained for different amounts of time, such as during 1 month or 1 year. We estimated Z values for C ⫽ 0.1 (10% mass loss) and t ⫽ 2 min for the PI1–5/0, PI1–5/20, and PI1–5/50 samples. Figure 10 shows plots of the predicted thermal longevity versus the degradation temperature calculated in such a way. These results predict sharply improved long-term thermal stability of the hybrids increasing with silica content. Therefore, for pure PI1, a 10% mass loss occurs for a few minutes at about 500 °C, whereas the same mass loss can presumably be attained after treatment for 1 year at about 200 °C. Meanwhile, for the PI1–5/20 and PI1–5/50 hybrids, such a degree of thermal degradation is observed for a short time at 520 –550 °C, but it may be expected as a result of a 1-year treatment at 320 and 380 °C, respectively. In other words, incorporating the silica nanophases into PI by means of hybridization presumably allows us to expect a large increase, of more than 100 °C, in an admissible working temperature in some cases. The thermal longevity of high-performance PI at 400 °C may by increased then by 3 orders of magnitude (Fig. 10).

(2)

where C ⫽ Wd/W0 is a fraction of a degraded, evaporated substance; Wd is the absolute mass of the evaporated substance; W0 is the initial sample

Figure 10. Predicted thermal longevity versus temperature dependence calculated from the thermogravimetric data obtained for PI1 and two indicated hybrids.

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CONCLUSIONS For the first time, molecular dynamics in nanostructured organic–inorganic polymeric hybrids prepared via a sol– gel process were studied in detail under different conditions in combination with a few complementary techniques. Silica nanoclusters, chemically bound with the ends of PI chains, may substantially modify molecular dynamics in this high-performance polymer. The effects of both large enhancement and drastic suppression of molecular motion were observed, depending on such factors as the hybrid composition, the PI chemical structure and chain length, and the temperature region of relaxation. The common interpretation of all the effects found was offered, which was based on the notions of (1) the peculiar dynamics in nanoscale spaces with three-dimensional confinement, (2) the influence of anchoring chains to rigid constraints on their dynamic behavior, and (3) the decisive role of Kuhn statistical segments in both the glass transition and ␤ relaxation in flexiblechain polymers. The contributions of three kinds of movements, associated with loosened molecular packing or suppressed dynamics (immobilizing factor) or with normal bulklike dynamic behavior, could be distinguished, to a certain extent, in PI–silica hybrids. Complete domination of the immobilizing factor in their segmental dynamics should be expected when the ratio of the contour chain length to the Kuhn segment length (l/A) is less than or equal to 3–5. Modifying a polymer (PI) by means of hybridization with inorganic nanoclusters may lead to a large increase in long-term thermal stability with respect to pure PI. The authors are grateful to the Russian Fund for Basic Research (grant 00-03-33173), International Association for the Promotion of Cooperation with Scientists from New Independent States of the Former Soviet Union (project 97-1936), NATO (grant PST.EV.97310 to V. A. Bershtein), the Greek Ministry of Development, and the Czech Ministry of Education, Youth and Sport for their financial support of this work.

REFERENCES AND NOTES 1. Novak, B. M. Adv Mater 1993, 5, 422. 2. Hybrid Organic–Inorganic Composites; Mark, J. E.; Lee, C. Y.-C.; Bianconi, P. A., Eds.; ACS Symposium Series 585; American Chemical Society: Washington, DC, 1995.

3. Lee, J. H.; Guo, L.; Beaucage, G.; Macip-Boulis, M. A.; Yang, A. J. M. J Polym Sci Part B: Polym Phys 1996, 34, 3073. 4. Noell, J. L. W.; Wilkes, D. L.; Mohanty, D. K.; McGrath, J. E. J Appl Polym Sci 1990, 40, 1177. 5. Mark, J. E.; Wang, S.; Ahmad, Z. Macromol Chem Macromol Symp 1995, 98, 701. 6. Zhou, W.; Dong, J. H.; Qiu, K. Y.; Wei, Y. J Polym Sci Part A: Polym Chem 1998, 36, 1607. 7. Huang, Z. H.; Qiu, K. Y.; Wei, Y. J Polym Sci Part A: Polym Chem 1997, 35, 2403. 8. Tian, D.; Dubois, P. H.; Jerome, K. J Polym Sci Part A: Polym Chem 1997, 35, 2295. 9. Wang, S.; Ahmad, Z.; Mark, J. E. Chem Mater 1994, 6, 943. 10. Morikawa, A.; Iyoki, Y.; Kakimoto, M.; Imai, Y. J Mater Chem 1992, 2, 679. 11. Morikawa, A.; Yamaguchi, H.; Kakimoto, M.; Imai, Y. Chem Mater 1994, 6, 913. 12. Sysel, P.; Pulec, R. Presented at the 4th European Technical Symposium on Polyimides and High Performance Polymers, Montpellier, France, 1996; PI 3–5. 13. Sysel, P.; Pulec, R.; Maryska, M. Polym J 1997, 29, 607. 14. Suizdak, D. A.; Mauritz, K. A. J Polym Sci Part B: Polym Phys 1999, 37, 143. 15. Hu, Q.; Marand, E. Polymer 1999, 40, 4833. 16. Mascia, L.; Kioul, A. J Mater Sci Lett 1994, 13, 641. 17. Iyoki, Y.; Kakimoto, M.; Imai, Y. High Perform Polym 1994, 6, 43. 18. Tamaki, R.; Naka, K.; Chujo, Y. Polym J 1998, 30, 60. 19. Matejka, L.; Dusek, K.; Plestil, J.; Kriz, J.; Lednicky, F. Polymer 1998, 40, 171. 20. Norisuye, T.; Shibayama, M.; Tamaki, R.; Chujo, Y. Macromolecules 1999, 32, 1528. 21. Smaihi, M.; Jermoumi, T.; Marigan, J.; Nobble, R. D. J Membr Sci 1996, 116, 211. 22. Moaddeb, M.; Koros, W. J. J Membr Sci 1997, 125, 143. 23. Yoshida, M.; Prassad, P. N. Appl Opt 1996, 35, 1500. 24. Krug, H.; Schmidt, H. New J Chem 1994, 18, 1125. 25. Dave, B. C.; Dunn, B.; Valentine, J. S.; Zink, J. I. Anal Chem 1994, 66, 1120. 26. Bershtein, V. A.; Egorova, L. M.; Yakushev, P. N.; Georgoussis, G.; Kyritsis, A.; Pissis, P.; Sysel, P.; Brozova, L. Macromol Chem Macromol Symp 1999, 146, 9. 27. Bershtein, V. A. Mechanohydrolytic Processes and Strength of Solids (in Russian); Nauka: Leningrad, 1987. 28. Bershtein, V. A.; Egorov, V. M. Differential Scanning Calorimetry of Polymers: Physics, Chemistry, Analysis, Technology; Ellis Horwood: New York, 1994.


29. Peschanskaya, N. N.; Yakushev, P. N.; Sinani, A. B.; Bershtein, V. A. Thermochim Acta (Therm Anal Calorim Polym Phys) 1994, 238, 429. 30. Peschanskaya, N. N.; Yakushev, P. N.; Sinani, A. B.; Bershtein, V. A. Macromol Chem Macromol Symp 1997, 119, 79. 31. Bershtein, V. A.; Yakushev, P. N.; Karabanova, L.; Sergeeva, L.; Pissis, P. J Polym Sci Part B: Polym Phys 1999, 37, 429. 32. Van Turnhout, J. In Electrets; Sessler, G. M., Ed.; Topics in Applied Physics 33; Springer-Verlag: Berlin, 1980; p 81. 33. Pissis, P.; Anagnostopoulou-Konsta, A.; Apekis, L.; Daoukaki-Diamanti, D.; Christodoulides, C. J NonCryst Solids 1991, 131–133, 1174. 34. Bessonov, M. I.; Kuznetsov, N. P.; Adrova, F. S.; Phlorinsky, F. S. Vysokomol Soedin 1974, 16, 2093. 35. Stauga, R.; Schoenhals, A.; Carius, H.-E.; Mudrak, C. V.; Privalko, V. P. New Polym Mater 1998, 5, 119. 36. Boiteux, G.; Oraison, J. M.; Seytre, G.; Rabilloud, G.; Sillion, B. In Polyimides and Other High-Temperature Polymers; Abadie, M. J. M.; Sillion, B., Eds.; Elsevier: Amsterdam, 1991. 37. Bershtein, V. A.; Egorov, V. M.; Emelyanov, Yu. A. Vysokomol Soedin A 1985, 27, 2440, 2451. 38. Bershtein, V. A.; Egorov, V. M.; Egorova, L. M.; Ryzhov, V. A. Thermochim Acta (Therm Anal Calorim Polym Phys) 1994, 238, 41. 39. Bershtein, V. A.; Ryzhov, V. A. Adv Polym Sci 1994, 114, 43. 40. Jackson, C. L.; McKenna, G. B. J Non-Cryst Solids 1991, 131–133, 221. 41. McKenna, G. B.; Jackson, C. L. Chem Mater 1996, 8, 2128. 42. Pissis, P.; Daoukaki-Diamanti, D.; Apekis, L.; Christodoulides, C. J Phys: Condens Matter 1994, 6, L325. 43. Keddie, J. L.; Jones, R. A. L.; Cory, R. A. Faraday Discuss 1994, 98, 219. 44. Prucker, O.; Christian, S.; Bock, H.; Frank, C. N.; Knoll, W. Polymer Preprints 1997, 88, 918. 45. Van Zanten, J. H.; Wallace, W. E.; Wu, W. L. Physical Review E 1996, 53, 2053.

1069

46. Mel’nichenko, Ju. B.; Schuller, J.; Richert, R.; Ewen, B. J Chem Phys 1995, 103, 2016. 47. Barut, G.; Pissis, P.; Pelster, R.; Nimtz, G. Phys Rev Lett 1998, 80, 3543. 48. Daoukaki, D.; Barut, G.; Pelster, R.; Nimtz, G.; Kyritsis, A.; Pissis, P. Physical Review B 1998, 58, 5336. 49. Gorbatschov, W.; Arndt, A.; Stannarus, R.; Kremer, F. Europhys Lett 1996, 35, 719. 50. DeMaggio, G. B.; Frieze, W. E.; Gidley, D. W.; Minz, Z.; Hristov, H. A.; Yee, A. F. Phys Rev Lett 1997, 78, 1524. 51. Forrest, J. A.; Dalnoki-Veress, K.; Stevens, J. R.; Dutcher, J. R. Phys Rev Lett 1996, 77, 2002. 52. Huwe, A.; Kremer, F.; Behrens, P.; Schweiger, W. Phys Rev Lett 1999, 82, 2338. 53. Vieweg, S.; Unger, R.; Hempel, E.; Donth, E. J Non-Cryst Solids 1998, 235–237, 470. 54. Gotlib, Yu. Ya.; Torchinsky, I. A.; Shevelev, V. A. Vysokomol Soedin A 1996, 38, 1966. 55. Bershtein, V. A.; Levin, V. Yu.; Egorova, L. M.; Egorov, V. M.; Zhdanov, A. A.; Slonimsky, G. L.; Rabkina, A. Yu.; Zavin, B. G.; Grischenko, O. I. Vysokomol Soedin A 1987, 29, 2360. 56. Bershtein, V. A.; Levin, V. Yu.; Egorova, L. M.; Egorov, V. M.; Zhdanov, A. A.; Slonimsky, G. L.; Makarova, L. I.; Martirosov, V. I. Vysokomol Soedin A 1987, 29, 2553. 57. Bershtein, V. A.; Egorova, L. M.; Sysel, P. J Macromol Sci Phys 1998, 37, 747. 58. Bershtein, V. A.; Egorov, V. M.; Egorova, L. M.; Sysel, P.; Zgonnik, V. N. J Non-Cryst Solids 1998, 235–237, 476. 59. Bershtein, V. A.; Egorov, V. M.; Podolsky, A. F.; Stepanov, V. A. J Polym Sci Polym Lett Ed 1985, 23, 371. 60. Mac Callum, J. R. In: Comprehensive Polymer Science; Allen, G.; Bevington, J. C., Eds.; Pergamon: Oxford, 1989; p 903. 61. Emanuel, N. M.; Buchachenko, A. L. Chemical Physics of Ageing and Stabilization of Polymers (in Russian); Nauka: Moscow, 1982. 62. Shi, L. T. Thermochim Acta 1990, 166, 127.