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Study of ethylene copolymers films by dielectric spectroscopy: influence of the polymer thickness on the glass-relaxation temperature
Progress in Organic Coatings 37 (1999) 49±56
Study of ethylene copolymers ®lms by dielectric spectroscopy: in¯uence of the polymer thickness on the glass-relaxation temperature S. Bistac*, M.F. Vallat, J. Schultz Institut de Chimie des Surfaces et Interfaces, ICSI-CNRS, 15, rue Jean Starcky, 68057 Mulhouse Cedex, France Received 4 February 1999; received in revised form 2 August 1999; accepted 4 August 1999
Abstract The aim of this work is to investigate the in¯uence of the thickness of a polymer layer deposited on a glass cover on the glass-relaxation temperature. The chosen polymer is an ethylene-vinylacetate copolymer (EVA) with different amounts of vinylacetate (VA). Dielectric spectroscopy analysis is performed directly on polymer±glass assemblies and the relaxation temperature corresponding to the maximum of the loss factor, tan , is examined. This temperature is related to the glass transition of the copolymer. The results show that at high thickness values, the glass-relaxation temperature corresponds to the glass-transition temperature of the bulk copolymer. However, when the polymer thickness decreases, the relaxation temperature increases greatly whatever is the VA content. These results evidence the formation of an interphase in the polymer layer, localised in the vicinity of the glass surface, and whose properties are different from that of the bulk. The increase of the relaxation temperature when the polymer thickness decreases is a consequence of the reduced mobility of the polymer chains in the interphasial zone. # 1999 Elsevier Science S.A. All rights reserved. Keywords: Dielectric spectroscopy; Interphase; Glass relaxation; EVA copolymers
1. Introduction During a dielectric experiment, a periodic potential is applied to the sample, and the capacitive and conductive responses of the sample are measured [1]. The dielectric constant * is a complex quantity and the ratio tan /0 is called the loss factor. Different polarisation processes can be responsible for a dielectric absorption peak. The atomic and electronic polarisations occur at elevated frequencies (109±1015 Hz), respectively, infrared and optical regions. At lower frequencies, in the radio range (typically between 102 and 107 Hz), the orientation of the dipoles are responsible to the dielectric absorption. At the lowest frequencies (low audio range) can appear the interfacial polarisation. In the presence of dipolar constituents, one or more absorption regions will be present, not all of them necessarily associated with the dipolar dispersions. At the lowest frequencies, signi®cantly large 00 values will arise from the d.c. conductivity of the medium, further electrode polarisa*
Corresponding author. Tel.: 00333-89-60-87-00; fax: 00333-89-60-8799 E-mail address: [email protected] (S. Bistac)
tion may well appear and interfacial polarisation will be found if the system is not homogeneous [2]. The characteristic dielectric properties of a polymer are largely determined by the chemical nature of the polymeric repeat unit. Complications may arise from the presence of second components such as plasticisers, impurities and adsorbed water or residual catalyst. Ionic d.c. conduction can be responsible for a strong rise of loss but only at low frequencies (and high temperatures). The presence of a great conductance is also liable to cause electrode polarisation effects at the lowest frequencies (ions in the immediate neighbourhood migrate to the electrode surfaces) [2,3]. If the dielectric specimen is not homogeneous, but contains regions of different permittivity, such as those that may arise from cracks in a solid specimen, the presence of inclusions (bubbles, etc.) in a composite (containing ®bres) or in a ®lled polymer, misleading values of permittivity may be found experimentally, due to a surface or interfacial polarisation mechanism. This dielectric effect is referred to as the Maxwell±Wagner effect. If the applied frequency is ranging from typically 102 to 7 10 Hz and if the analysed sample is homogeneous, the dielectric absorption can be truly ascribable to a dipole reorientation process. Polymeric materials are characterised
0300-9440/99/$ ± see front matter # 1999 Elsevier Science S.A. All rights reserved. PII: S 0 3 0 0 - 9 4 4 0 ( 9 8 ) 0 0 0 4 8 - 4
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S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56
by the presence of a phase transition at the so-called glasstransition temperature Tg. At this temperature appear local segmental motions (relaxation of small portions of the backbone chain). These rearrangements involve cooperative thermal motions of individual chain segments. At Tg, there is a sudden change in many properties, e.g., heat capacity, modulus, thermal expansion, and the dielectric behaviour changes likewise, especially if the polymer is polar. At Tg, a loss peak is present, due to the energy dissipation in the polymer induced by the dipoles reorientation. Relaxation arises from the motion of chains to a new equilibrium distribution of conformations subsequent to the application of an external stress or ®eld. Frequently, in polymer, the glass±rubber transition is mechanically and dielectrically the most prominent one. In some cases, other relaxation processes can be observed below Tg (subglass relaxations). These secondary relaxations result from motions within the polymer in the glass-like state (rotation of side-group in particular) [4]. In the literature, viscoelastic and dielectric properties of various polymers have been studied in a systematic and comparative way, that reinforce the evidence of a direct relation (in some particular experimental conditions) between the dielectric loss of a polymer and its glasstransition relaxation [5±8]. Deutsch et al. [6] have studied the relation between the structure of polymers and their dynamic mechanical and dielectrical properties. These two properties were found to be intrinsically correlated. Each of the mechanical and dielectrical dispersion regions found in a given polymer is associated with a de®nable structural feature, i.e., groups of atoms in the polymer, such that each group gives rise to a mechanical dispersion in the same temperature region and, if polar, also to a dielectric dispersion in the same temperature and frequency regions [6]. Cormenero et al. [8] have investigated the correspondence between dielectric and mechanical relaxations in poly(vinyl ethylene). They conclude that the shift factors, describing the temperature dependence of the segmental relaxation time, are equivalent when measured dielectrically or mechanically [8]. In the glass-transition region, the loss factor goes through a maximum and dynamic testing appears then to be a very sensitive method of determining the glass transition of polymers. One of the main advantages of dielectric spectroscopy, compared to the dynamic mechanical measurements, is the great sample preparation facilities. Thin polymer ®lms deposited on a substrate can be more easily studied by dielectric spectroscopy and no ``mechanical coupling effect'', i.e., contribution and in¯uence of the mechanical response of the substrate, appears [9]. The properties of a polymer in close contact with a substrate can be sometimes different from that of the bulk polymer. The interphase represents the region localised near the interface, where the chemical, mechanical or crystalline
properties, for example, are different from ones in the bulk. Modi®cations of the polymer properties can have various origins. The curing process is able to induce modi®cations of the polymer properties in the vicinity of the substrate [10,11]. The crystallisation kinetic and morphology of thermoplastic polymers can be affected by the contact of a substrate [12,13]. Modi®cations of polymer properties can also be induced by simple adsorption of macromolecules onto a surface [14]. In the present work, the dielectric properties of thin ®lms of ethylene-vinylacetate (EVA) copolymers deposited on a glass substrate are studied. EVA copolymers are polar materials, the dipole being largely in the carbonyl group of the acetate group [2]. Different polymer thicknesses are obtained and the in¯uence on the dielectric response is considered. By decreasing the polymer thickness, the relative proportion of a possible interphase increases and the properties of this interphasial zone become therefore detectable. 2. Experimental Three copolymers of ethylene-vinylacetate (EVA) with various amounts of vinylacetate (VA) Ð 14, 28 and 40 wt% Ð are studied. EVA copolymers are purchased from Elf Atochem. Their degree of crystallinity is evaluated by differential scanning calorimetry (DSC) with a scan rate of 108C/min from ÿ1008C to 1508C. For ®lm preparation, EVA at 28%VA and 40%VA are dissolved in trichloroethylene and polymer solutions, with a concentration of 1 mol/l, and cast on clean cover glasses (35 mm 35 mm, thickness 1 mm). Different thicknesses (from 1 to 30 mm) are obtained by varying the number of deposited layers. These samples are then heated at 1508C during 6 min in order to eliminate the residual solvent and allow the relaxation of the macromolecules. EVA at 14%VA, which is more crystalline, is insoluble in trichloroethylene (and in other usual solvents) and is therefore moulded on glass plates in heating press, at 1508C during 6 min. Thin thicknesses cannot be obtained using this ``hot melt'' technique (minimum thickness 30 mm). Thicker samples of each EVA are obtained by the same technique, they can be used as reference bulk samples. Dielectric experiments are carried out with a DETA instrument from Rheometric Scienti®c. Both thin polymer ®lms deposited on glass plates and thicker bulk EVA ®lms, inserted between two metallic plates, are analysed at 10 kHz, from ÿ608C to 508C at a 28/min scanning rate. The uncovered glass plate is also analysed in order to verify the absence of any relaxation peak, in the studied range of frequency and temperatures. 3. Results and discussion Table 1 presents the DSC results obtained for the bulk EVA. The degree of crystallinity is calculated by using the
S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56 Table 1 Melting enthalpy and degree of crystallinity of EVA copolymers
EVA 14%VA EVA 28%VA EVA 40%VA
Melting enthalpy H (mJ/g)
Degree of crystallinity (%)
97 34 13
35 12 5
melting enthalpy of pure polyethylene (280 J/mol [15]), assuming that only polyethylene sequences crystallise [16]. The results indicate that the degree of crystallinity of the studied EVA depends on the vinylacetate content. A glass-transition temperature close to ÿ208C is detected for all the EVA samples, whatever the VA content. This result is in agreement with the value found in the literature [17]. Figs. 1±6 show the tan curves obtained by DETA for the different EVA. At this frequency measurement (10 kHz), atomic and electronic polarisations are not detectable. At 10 kHz, and in the temperature range of the tan peak, a conduction process can be excluded. Moreover, the studied EVA copolymers can be considered as homogeneous materials. No charges, no ®bres and no inclusions are present. The interfacial polarisation can also be neglected. For EVA copolymers, a subglass relaxation can be observed, but at lower temperature (around ÿ1008C) and the corresponding peak is only slightly pronounced [18]. The temperature at which tan goes through a maximum can be related, in these experimental conditions, to the glasstransition temperature, Tg. Fig. 1 shows the tan curve obtained by DETA for bulk EVA 14%VA. For a bulk ®lm (450 mm thick), the tempera-
51
ture of the tan peak is equal to ÿ78C. Similar curves are obtained for the other EVA. Figs. 2 and 3 present the tan curves of bulk EVA 28%VA and 40%VA, respectively. The maximum of the tan peak appears in the same range of temperature (ÿ58C; ÿ108C). These values are in agreement with the values found in the literature [4]. Indeed, Buerger and Boyd. [18] have studied the dielectric relaxation processes in ethylene vinylacetate copolymers and have attributed the tan peak located around ÿ58C at 10 kHz to the glass-transition process. The temperature measured for these bulk EVA ®lms are considered as a reference for the study of the thinner EVA ®lms DETA measurements are then performed on samples with different polymer thicknesses deposited on glass. Fig. 4 shows the tan curves obtained for EVA 14%VA. It appears that the tan peak is greatly affected in both magnitude and position when the polymer thickness varies. The intensity of the damping peak decreases when the polymer thickness is reduced, due to the fact that the dissipative layer (EVA) is less important. But, more interesting, the temperature of the glass relaxation is shifted signi®cantly towards higher temperature when the thickness of the polymer ®lm decreases. For higher thicknesses, the value of the glass-relaxation temperature reaches that obtained for the bulk EVA ®lm. These phenomena are also observed for EVA 28%VA and 40%VA as seen on Figs. 5 and 6. Furthermore, the tan peak is not symmetrical for the smaller thickness, with a broadening towards a temperature close to the bulk EVA glassrelaxation temperature. Table 2 summarises the results obtained for the different EVA ®lms.
Fig. 1. tan curve of bulk EVA 14%VA.
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S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56
Fig. 2. tan curve of bulk EVA 28%VA.
Fig. 3. tan curve of bulk EVA 40%VA. Table 2 Temperature of the maximum of the tan peak obtained for the EVA 14%VA, 28%VA and 40%VA at different thicknesses (T8 18C) EVA 14%VA
Thickness (mm) T8 tan peak (8C)
Bulk sample ÿ7.4
450 ÿ7.3
70 7
30 6
EVA 28%VA
Thickness (mm) T8 tan peak (8C)
Bulk sample ÿ10
30 ÿ10
10 ÿ5.3
1 10.6
EVA 40%VA
Thickness (mm) T8 tan peak (8C)
Bulk sample ÿ5
10 ÿ5.2
5 3.6
1 7
S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56
53
Fig. 4. tan curves of EVA 14%VA obtained for different polymer thicknesses.
Fig. 5. tan curves of EVA 28%VA obtained for different polymer thicknesses.
The results obtained by dielectric spectroscopy show therefore that the tan peak is shifted towards higher temperatures when the thickness of the polymer ®lm decreases, whatever is the VA content. At higher polymer thickness, the value of the glass-relaxation temperature corresponds to the value obtained for the bulk EVA ®lms. At a molecular scale, the glass-relaxation temperature re¯ects the mobility of the macromolecules in the ®lms.
An increase of this temperature, as observed for thin polymer layers, corresponds therefore to a decrease of the mobility of the chains. This reduction in segmental mobility at lower polymer thickness indicates that the reduction of mobility probably affects a region located in the vicinity of the polymer±substrate interface. This result suggests the formation of an interphase, that exhibit different viscoelastic properties compared to the bulk ones. The increase of the
54
S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56
Fig. 6. tan curves of EVA 40%VA obtained for different polymer thicknesses.
glass-relaxation temperature is already reported in the literature for polymer±substrate assemblies. Also such an increase in the glass-transition temperature is observed by dynamic mechanical measurements in the case of vinylacetate copolymers introduced in a steel±polymer±steel sandwich [19,20]: the temperature of the tan peak, attributed to the glass-transition relaxation, increases greatly when the EVA thickness decreases. For example, the temperature of the tan peak is shifted from ÿ228C (for a thick ®lm of EVA 28%VA) to 98C (for a thickness equal to 50 mm) [20]. These observations are comparable to our dielectric spectroscopy results. This comparison of the results obtained using both techniques, mechanical and dielectric spectroscopies, shows clearly that the observed dielectric relaxation is directly correlated to the glass-transition relaxation, and that the shift of the dielectric tan peak towards the higher temperatures, corresponds therefore to an increase of the glass-transition temperature, and not to a further relaxation process, at least in our experimental conditions (frequency, temperature, homogeneity of the sample, etc.). Results obtained for the other systems indicate various effects of the substrate. A study of epoxy-based assemblies has also shown an increase of the glass-relaxation temperature when the polymer thickness decreases, attributed to the presence of an interphase of overcrosslinked polymer in contact with the substrate [21]. Studies of PMMA layers deposited on a substrate have also evidenced differences in the mobility of the chains when the polymer thickness varies [22]. The observed phenomena for EVA copolymers can therefore be explained by the formation of an interphasial layer
close to the interface in which the amorphous phase is constrained. For thin polymer ®lms, a broadening of the tan peak toward a temperature close to the bulk EVA glassrelaxation temperature corresponds doubtless to chains whose mobility is not affected by the interface and then localised, in principle, far from the interface. The tan spectra do not exhibit only single relaxation peak, especially when the EVA layer is thin. The broadening of the tan peak observed for the lower polymer thickness is due to the increase of the distribution of the relaxation time. When the polymer thickness decreases, the mobility of some chains (present in the interphasial region) is reduced. However, some macromolecules, preferentially located far from the interface and not in the interphasial region, are not perturbed and their mobility corresponds to the mobility of the bulk polymer chains (Tg equal to the Tg measured for the thicker ®lms). Moreover, a probable gradient of mobility exists from the ``core'' of the ®lm (non-perturbed chains) to the interphasial zone (constrained chains), that leads to a large distribution of the relaxations times, explaining the broadening of the tan peak. Such a broadening is also observed, for the lower EVA thickness, on the tan peaks obtained by dynamic mechanical measurements [20]. The peak is also non-symmetrical. This effect can be due to the different contribution of the various relaxation times. For the lowest EVA thickness, the contribution of the high temperature component is higher compared to the other thicknesses, indicating that the constrained chains are more numerous compared to the non-perturbed chains. It can be speci®ed that the presented spectra are obtained from true experimental values, directly given by the appa-
S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56
ratus. Any corrections are made, that explains the fact that the baseline of the spectra is not perfect. Elsewhere, it appears important to say that the tan peak shift is not affected by the sample preparation. No difference can be observed between the compression molding way (heating press used for the EVA 14%VA) and the solution casting way (used for EVA 28%VA and 40%VA). This result seems to indicate that no con®nement effects (due to the pressure applied to the polymer ®lm during the melt process) occur, a variation of the glass-transition temperature being also observed for the casting way, without any mechanical strain. The reduced mobility of the polymer chains in the vicinity of the interface with the substrate can be explained by a change of the crystalline organisation in a layer close to the interface. Indeed, previous studies have evidenced some change of the crystalline organisation in thin EVA ®lms analysed by differential scanning calorimetry [20]. The results have shown that more numerous small and disorganised crystals are present in the vicinity of the interface. These crystals act as physical ties that reduce the mobility of the amorphous chains which link them. This crystalline modi®cation could explain the reduction of the mobility at lower EVA thickness. Moreover, for interfacial energy minimisation reasons, an orientation of the vinylacetate groups towards the glass substrate doubtless occurs. These migration and orientation phenomena of the polar groups of the EVA are favoured during the heating step (under the heating press for the EVA 14% or into the oven for the solvent cast ®lms) and can induce the formation of an interphasial layer close to the interface with the substrate. Considering now the effect of the VA content on the increase of the glass transition at low thickness, an important increase of the Tg is observed for the EVA 14%VA for thickness about 70 mm. For EVA 28%VA and 40%VA, no increase of Tg is seen at the same thickness range. The formation of an interphase of reduced mobility also appears to be favoured in the case of EVA 14%VA. A temperature of about 78C is reached for an EVA 14%VA, 70 mm thick and for a thickness close to only 1 mm for both EVA 28%VA and 40%VA. The greater increase of Tg for EVA 14%VA could be explained by the fact that this EVA is the most crystalline. The mobility of the amorphous chains will be consequently more sensitive to the least modi®cation of the crystallinity organisation. Furthermore, it is interesting to notice that an identical value of Tg, close to 7±108C is reached at low thickness for all the EVA. Indeed, despite the fact that no layer thinner than 30 mm can be obtained for EVA 14%VA, the glass-transition temperature seems to stabilise at a maximum close to 88C, no difference in Tg being observed between the 70 and 30 mm thick ®lms. These results indicate that the minimum mobility of the interphase chains is reached and is independent of the VA content of the EVA. But nevertheless, some additional studies on thinner ®lms are necessary to con®rm the stabilisation of the Tg value for lower thicknesses.
55
4. Conclusion The study of the glass-transition temperature of EVA ®lms deposited on a glass substrate by dielectric spectroscopy (DETA) leads to the following conclusions: In the experimental conditions used in this study, the observed tan peak corresponds to the glass-transition relaxation. The temperature corresponding to the maximum of the peak is related to the glass-transition temperature, Tg. The results obtained by DETA show that the glasstransition temperature is significantly shifted towards higher values when the polymer thickness is decreased, whatever is the VA content. These results indicate that the segmental mobility of the polymer chains is reduced at lower thicknesses, that corresponds to a mobility reduction in the vicinity of the interface with the substrate. These phenomena characterise the existence of a constrained interphase, localised close to the substrate and where the polymer properties differ from the bulk ones. Influence of a confinement effect due to mechanical strains can be rejected. No major effects of the compression process under heating press (used for the EVA 14%VA) or the solvent casting way (used for EVA 28%VA and 40%VA) have been noticed: both processes lead to films presenting an increase of Tg at lower thickness. Finally, dielectric spectroscopy appears to be a easier, sensitive and fruitful technique for studying relaxation phenomena in thin polymer films deposited on a substrate. References [1] K. Varadarajan, J. Coat. Technol. 55704 (1983) 95. [2] N.E. Hill, W.E. Vaughan, A.H. Price, M. Davies, Dielectric Properties and Molecular Behaviour, Van Nostrand Reinhold, London, 1969. [3] A.K. Jonscher, Dielectric Relaxation in Solids, Chelsea Dielectric Press, London, 1983. [4] N.G. Mc Crum, B.E. Read, G. Williams, in: Anelastic and Dielectric Effects in Polymeric Solids, Dover, New York, 1991. [5] O. Pelissou, R. Diaz Calleja, L. Gargallo, D. Radic, Polymer 3516 (1994) 3449. [6] K. Deutsch, E.A.W. Hoff, W. Reddish, J. Polym. Sci. 13 (1954) 565. [7] K.L. Ngai, S. Mashimo, G. Fytas, Macromolecules 21 (1988) 3030. [8] J. Cormenero, A. Alegria, P.G. Santangelo, K.L. Ngai, C.M. Roland, Macromolecules 27 (1994) 407. [9] J.V. Standish, H. LeidheiserJr., J. Coat. Technol. 53678 (1981) 53. [10] R.G. Dillingham, F.J. Boerio, J. Adhesion 24 (1987) 315. [11] J.S. Crompton, J. Mater. Sci. 24 (1989) 1575. [12] H. Schonhorn, J. Polym. Sci. 5 (1967) 919. [13] M. Nardin, E.M. Asloun, F. Muller, J. Schultz, Polym. Adv. Technol. 2 (1991) 161. [14] C.A. Kumins, C.J. Knauss, R.J. Ruch, Progress Org. Coatings 28 (1996) 17.
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[15] J. Brandrup, E.H. Immergut, Polymer Handbook, 2nd ed., Wiley, New York, 1975. [16] R.J. Koopmans, R. Van Der Linden, E.F. Vansant, J. Adhesion 11 (1980) 191. [17] L.E. Nielsen, J. Polym. Sci. 13 (1960) 357. [18] D.E. Buerger, R.H. Boyd, Macromolecules 22 (1989) 2699. [19] H. Braunisch, Acustica 22 (1969-1970) 136.
[20] S. Bistac, D. Cheret, M.F. Vallat, J. Schultz, J. Applied Polym. Sci. 65 (1997) 347. [21] H. Bourhala, J. Chauchard, J. Lenoir, H. Romand, Die Angewandte Makromolecular Chemie 1782941 (1990) 47. [22] O.N. Tretinnikov, R.G. Zhabankov, J. Mater. Sci. Lett. 10 (1991) 1032.
Study of ethylene copolymers ®lms by dielectric spectroscopy: in¯uence of the polymer thickness on the glass-relaxation temperature S. Bistac*, M.F. Vallat, J. Schultz Institut de Chimie des Surfaces et Interfaces, ICSI-CNRS, 15, rue Jean Starcky, 68057 Mulhouse Cedex, France Received 4 February 1999; received in revised form 2 August 1999; accepted 4 August 1999
Abstract The aim of this work is to investigate the in¯uence of the thickness of a polymer layer deposited on a glass cover on the glass-relaxation temperature. The chosen polymer is an ethylene-vinylacetate copolymer (EVA) with different amounts of vinylacetate (VA). Dielectric spectroscopy analysis is performed directly on polymer±glass assemblies and the relaxation temperature corresponding to the maximum of the loss factor, tan , is examined. This temperature is related to the glass transition of the copolymer. The results show that at high thickness values, the glass-relaxation temperature corresponds to the glass-transition temperature of the bulk copolymer. However, when the polymer thickness decreases, the relaxation temperature increases greatly whatever is the VA content. These results evidence the formation of an interphase in the polymer layer, localised in the vicinity of the glass surface, and whose properties are different from that of the bulk. The increase of the relaxation temperature when the polymer thickness decreases is a consequence of the reduced mobility of the polymer chains in the interphasial zone. # 1999 Elsevier Science S.A. All rights reserved. Keywords: Dielectric spectroscopy; Interphase; Glass relaxation; EVA copolymers
1. Introduction During a dielectric experiment, a periodic potential is applied to the sample, and the capacitive and conductive responses of the sample are measured [1]. The dielectric constant * is a complex quantity and the ratio tan /0 is called the loss factor. Different polarisation processes can be responsible for a dielectric absorption peak. The atomic and electronic polarisations occur at elevated frequencies (109±1015 Hz), respectively, infrared and optical regions. At lower frequencies, in the radio range (typically between 102 and 107 Hz), the orientation of the dipoles are responsible to the dielectric absorption. At the lowest frequencies (low audio range) can appear the interfacial polarisation. In the presence of dipolar constituents, one or more absorption regions will be present, not all of them necessarily associated with the dipolar dispersions. At the lowest frequencies, signi®cantly large 00 values will arise from the d.c. conductivity of the medium, further electrode polarisa*
Corresponding author. Tel.: 00333-89-60-87-00; fax: 00333-89-60-8799 E-mail address: [email protected] (S. Bistac)
tion may well appear and interfacial polarisation will be found if the system is not homogeneous [2]. The characteristic dielectric properties of a polymer are largely determined by the chemical nature of the polymeric repeat unit. Complications may arise from the presence of second components such as plasticisers, impurities and adsorbed water or residual catalyst. Ionic d.c. conduction can be responsible for a strong rise of loss but only at low frequencies (and high temperatures). The presence of a great conductance is also liable to cause electrode polarisation effects at the lowest frequencies (ions in the immediate neighbourhood migrate to the electrode surfaces) [2,3]. If the dielectric specimen is not homogeneous, but contains regions of different permittivity, such as those that may arise from cracks in a solid specimen, the presence of inclusions (bubbles, etc.) in a composite (containing ®bres) or in a ®lled polymer, misleading values of permittivity may be found experimentally, due to a surface or interfacial polarisation mechanism. This dielectric effect is referred to as the Maxwell±Wagner effect. If the applied frequency is ranging from typically 102 to 7 10 Hz and if the analysed sample is homogeneous, the dielectric absorption can be truly ascribable to a dipole reorientation process. Polymeric materials are characterised
0300-9440/99/$ ± see front matter # 1999 Elsevier Science S.A. All rights reserved. PII: S 0 3 0 0 - 9 4 4 0 ( 9 8 ) 0 0 0 4 8 - 4
50
S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56
by the presence of a phase transition at the so-called glasstransition temperature Tg. At this temperature appear local segmental motions (relaxation of small portions of the backbone chain). These rearrangements involve cooperative thermal motions of individual chain segments. At Tg, there is a sudden change in many properties, e.g., heat capacity, modulus, thermal expansion, and the dielectric behaviour changes likewise, especially if the polymer is polar. At Tg, a loss peak is present, due to the energy dissipation in the polymer induced by the dipoles reorientation. Relaxation arises from the motion of chains to a new equilibrium distribution of conformations subsequent to the application of an external stress or ®eld. Frequently, in polymer, the glass±rubber transition is mechanically and dielectrically the most prominent one. In some cases, other relaxation processes can be observed below Tg (subglass relaxations). These secondary relaxations result from motions within the polymer in the glass-like state (rotation of side-group in particular) [4]. In the literature, viscoelastic and dielectric properties of various polymers have been studied in a systematic and comparative way, that reinforce the evidence of a direct relation (in some particular experimental conditions) between the dielectric loss of a polymer and its glasstransition relaxation [5±8]. Deutsch et al. [6] have studied the relation between the structure of polymers and their dynamic mechanical and dielectrical properties. These two properties were found to be intrinsically correlated. Each of the mechanical and dielectrical dispersion regions found in a given polymer is associated with a de®nable structural feature, i.e., groups of atoms in the polymer, such that each group gives rise to a mechanical dispersion in the same temperature region and, if polar, also to a dielectric dispersion in the same temperature and frequency regions [6]. Cormenero et al. [8] have investigated the correspondence between dielectric and mechanical relaxations in poly(vinyl ethylene). They conclude that the shift factors, describing the temperature dependence of the segmental relaxation time, are equivalent when measured dielectrically or mechanically [8]. In the glass-transition region, the loss factor goes through a maximum and dynamic testing appears then to be a very sensitive method of determining the glass transition of polymers. One of the main advantages of dielectric spectroscopy, compared to the dynamic mechanical measurements, is the great sample preparation facilities. Thin polymer ®lms deposited on a substrate can be more easily studied by dielectric spectroscopy and no ``mechanical coupling effect'', i.e., contribution and in¯uence of the mechanical response of the substrate, appears [9]. The properties of a polymer in close contact with a substrate can be sometimes different from that of the bulk polymer. The interphase represents the region localised near the interface, where the chemical, mechanical or crystalline
properties, for example, are different from ones in the bulk. Modi®cations of the polymer properties can have various origins. The curing process is able to induce modi®cations of the polymer properties in the vicinity of the substrate [10,11]. The crystallisation kinetic and morphology of thermoplastic polymers can be affected by the contact of a substrate [12,13]. Modi®cations of polymer properties can also be induced by simple adsorption of macromolecules onto a surface [14]. In the present work, the dielectric properties of thin ®lms of ethylene-vinylacetate (EVA) copolymers deposited on a glass substrate are studied. EVA copolymers are polar materials, the dipole being largely in the carbonyl group of the acetate group [2]. Different polymer thicknesses are obtained and the in¯uence on the dielectric response is considered. By decreasing the polymer thickness, the relative proportion of a possible interphase increases and the properties of this interphasial zone become therefore detectable. 2. Experimental Three copolymers of ethylene-vinylacetate (EVA) with various amounts of vinylacetate (VA) Ð 14, 28 and 40 wt% Ð are studied. EVA copolymers are purchased from Elf Atochem. Their degree of crystallinity is evaluated by differential scanning calorimetry (DSC) with a scan rate of 108C/min from ÿ1008C to 1508C. For ®lm preparation, EVA at 28%VA and 40%VA are dissolved in trichloroethylene and polymer solutions, with a concentration of 1 mol/l, and cast on clean cover glasses (35 mm 35 mm, thickness 1 mm). Different thicknesses (from 1 to 30 mm) are obtained by varying the number of deposited layers. These samples are then heated at 1508C during 6 min in order to eliminate the residual solvent and allow the relaxation of the macromolecules. EVA at 14%VA, which is more crystalline, is insoluble in trichloroethylene (and in other usual solvents) and is therefore moulded on glass plates in heating press, at 1508C during 6 min. Thin thicknesses cannot be obtained using this ``hot melt'' technique (minimum thickness 30 mm). Thicker samples of each EVA are obtained by the same technique, they can be used as reference bulk samples. Dielectric experiments are carried out with a DETA instrument from Rheometric Scienti®c. Both thin polymer ®lms deposited on glass plates and thicker bulk EVA ®lms, inserted between two metallic plates, are analysed at 10 kHz, from ÿ608C to 508C at a 28/min scanning rate. The uncovered glass plate is also analysed in order to verify the absence of any relaxation peak, in the studied range of frequency and temperatures. 3. Results and discussion Table 1 presents the DSC results obtained for the bulk EVA. The degree of crystallinity is calculated by using the
S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56 Table 1 Melting enthalpy and degree of crystallinity of EVA copolymers
EVA 14%VA EVA 28%VA EVA 40%VA
Melting enthalpy H (mJ/g)
Degree of crystallinity (%)
97 34 13
35 12 5
melting enthalpy of pure polyethylene (280 J/mol [15]), assuming that only polyethylene sequences crystallise [16]. The results indicate that the degree of crystallinity of the studied EVA depends on the vinylacetate content. A glass-transition temperature close to ÿ208C is detected for all the EVA samples, whatever the VA content. This result is in agreement with the value found in the literature [17]. Figs. 1±6 show the tan curves obtained by DETA for the different EVA. At this frequency measurement (10 kHz), atomic and electronic polarisations are not detectable. At 10 kHz, and in the temperature range of the tan peak, a conduction process can be excluded. Moreover, the studied EVA copolymers can be considered as homogeneous materials. No charges, no ®bres and no inclusions are present. The interfacial polarisation can also be neglected. For EVA copolymers, a subglass relaxation can be observed, but at lower temperature (around ÿ1008C) and the corresponding peak is only slightly pronounced [18]. The temperature at which tan goes through a maximum can be related, in these experimental conditions, to the glasstransition temperature, Tg. Fig. 1 shows the tan curve obtained by DETA for bulk EVA 14%VA. For a bulk ®lm (450 mm thick), the tempera-
51
ture of the tan peak is equal to ÿ78C. Similar curves are obtained for the other EVA. Figs. 2 and 3 present the tan curves of bulk EVA 28%VA and 40%VA, respectively. The maximum of the tan peak appears in the same range of temperature (ÿ58C; ÿ108C). These values are in agreement with the values found in the literature [4]. Indeed, Buerger and Boyd. [18] have studied the dielectric relaxation processes in ethylene vinylacetate copolymers and have attributed the tan peak located around ÿ58C at 10 kHz to the glass-transition process. The temperature measured for these bulk EVA ®lms are considered as a reference for the study of the thinner EVA ®lms DETA measurements are then performed on samples with different polymer thicknesses deposited on glass. Fig. 4 shows the tan curves obtained for EVA 14%VA. It appears that the tan peak is greatly affected in both magnitude and position when the polymer thickness varies. The intensity of the damping peak decreases when the polymer thickness is reduced, due to the fact that the dissipative layer (EVA) is less important. But, more interesting, the temperature of the glass relaxation is shifted signi®cantly towards higher temperature when the thickness of the polymer ®lm decreases. For higher thicknesses, the value of the glass-relaxation temperature reaches that obtained for the bulk EVA ®lm. These phenomena are also observed for EVA 28%VA and 40%VA as seen on Figs. 5 and 6. Furthermore, the tan peak is not symmetrical for the smaller thickness, with a broadening towards a temperature close to the bulk EVA glassrelaxation temperature. Table 2 summarises the results obtained for the different EVA ®lms.
Fig. 1. tan curve of bulk EVA 14%VA.
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S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56
Fig. 2. tan curve of bulk EVA 28%VA.
Fig. 3. tan curve of bulk EVA 40%VA. Table 2 Temperature of the maximum of the tan peak obtained for the EVA 14%VA, 28%VA and 40%VA at different thicknesses (T8 18C) EVA 14%VA
Thickness (mm) T8 tan peak (8C)
Bulk sample ÿ7.4
450 ÿ7.3
70 7
30 6
EVA 28%VA
Thickness (mm) T8 tan peak (8C)
Bulk sample ÿ10
30 ÿ10
10 ÿ5.3
1 10.6
EVA 40%VA
Thickness (mm) T8 tan peak (8C)
Bulk sample ÿ5
10 ÿ5.2
5 3.6
1 7
S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56
53
Fig. 4. tan curves of EVA 14%VA obtained for different polymer thicknesses.
Fig. 5. tan curves of EVA 28%VA obtained for different polymer thicknesses.
The results obtained by dielectric spectroscopy show therefore that the tan peak is shifted towards higher temperatures when the thickness of the polymer ®lm decreases, whatever is the VA content. At higher polymer thickness, the value of the glass-relaxation temperature corresponds to the value obtained for the bulk EVA ®lms. At a molecular scale, the glass-relaxation temperature re¯ects the mobility of the macromolecules in the ®lms.
An increase of this temperature, as observed for thin polymer layers, corresponds therefore to a decrease of the mobility of the chains. This reduction in segmental mobility at lower polymer thickness indicates that the reduction of mobility probably affects a region located in the vicinity of the polymer±substrate interface. This result suggests the formation of an interphase, that exhibit different viscoelastic properties compared to the bulk ones. The increase of the
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S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56
Fig. 6. tan curves of EVA 40%VA obtained for different polymer thicknesses.
glass-relaxation temperature is already reported in the literature for polymer±substrate assemblies. Also such an increase in the glass-transition temperature is observed by dynamic mechanical measurements in the case of vinylacetate copolymers introduced in a steel±polymer±steel sandwich [19,20]: the temperature of the tan peak, attributed to the glass-transition relaxation, increases greatly when the EVA thickness decreases. For example, the temperature of the tan peak is shifted from ÿ228C (for a thick ®lm of EVA 28%VA) to 98C (for a thickness equal to 50 mm) [20]. These observations are comparable to our dielectric spectroscopy results. This comparison of the results obtained using both techniques, mechanical and dielectric spectroscopies, shows clearly that the observed dielectric relaxation is directly correlated to the glass-transition relaxation, and that the shift of the dielectric tan peak towards the higher temperatures, corresponds therefore to an increase of the glass-transition temperature, and not to a further relaxation process, at least in our experimental conditions (frequency, temperature, homogeneity of the sample, etc.). Results obtained for the other systems indicate various effects of the substrate. A study of epoxy-based assemblies has also shown an increase of the glass-relaxation temperature when the polymer thickness decreases, attributed to the presence of an interphase of overcrosslinked polymer in contact with the substrate [21]. Studies of PMMA layers deposited on a substrate have also evidenced differences in the mobility of the chains when the polymer thickness varies [22]. The observed phenomena for EVA copolymers can therefore be explained by the formation of an interphasial layer
close to the interface in which the amorphous phase is constrained. For thin polymer ®lms, a broadening of the tan peak toward a temperature close to the bulk EVA glassrelaxation temperature corresponds doubtless to chains whose mobility is not affected by the interface and then localised, in principle, far from the interface. The tan spectra do not exhibit only single relaxation peak, especially when the EVA layer is thin. The broadening of the tan peak observed for the lower polymer thickness is due to the increase of the distribution of the relaxation time. When the polymer thickness decreases, the mobility of some chains (present in the interphasial region) is reduced. However, some macromolecules, preferentially located far from the interface and not in the interphasial region, are not perturbed and their mobility corresponds to the mobility of the bulk polymer chains (Tg equal to the Tg measured for the thicker ®lms). Moreover, a probable gradient of mobility exists from the ``core'' of the ®lm (non-perturbed chains) to the interphasial zone (constrained chains), that leads to a large distribution of the relaxations times, explaining the broadening of the tan peak. Such a broadening is also observed, for the lower EVA thickness, on the tan peaks obtained by dynamic mechanical measurements [20]. The peak is also non-symmetrical. This effect can be due to the different contribution of the various relaxation times. For the lowest EVA thickness, the contribution of the high temperature component is higher compared to the other thicknesses, indicating that the constrained chains are more numerous compared to the non-perturbed chains. It can be speci®ed that the presented spectra are obtained from true experimental values, directly given by the appa-
S. Bistac et al. / Progress in Organic Coatings 37 (1999) 49±56
ratus. Any corrections are made, that explains the fact that the baseline of the spectra is not perfect. Elsewhere, it appears important to say that the tan peak shift is not affected by the sample preparation. No difference can be observed between the compression molding way (heating press used for the EVA 14%VA) and the solution casting way (used for EVA 28%VA and 40%VA). This result seems to indicate that no con®nement effects (due to the pressure applied to the polymer ®lm during the melt process) occur, a variation of the glass-transition temperature being also observed for the casting way, without any mechanical strain. The reduced mobility of the polymer chains in the vicinity of the interface with the substrate can be explained by a change of the crystalline organisation in a layer close to the interface. Indeed, previous studies have evidenced some change of the crystalline organisation in thin EVA ®lms analysed by differential scanning calorimetry [20]. The results have shown that more numerous small and disorganised crystals are present in the vicinity of the interface. These crystals act as physical ties that reduce the mobility of the amorphous chains which link them. This crystalline modi®cation could explain the reduction of the mobility at lower EVA thickness. Moreover, for interfacial energy minimisation reasons, an orientation of the vinylacetate groups towards the glass substrate doubtless occurs. These migration and orientation phenomena of the polar groups of the EVA are favoured during the heating step (under the heating press for the EVA 14% or into the oven for the solvent cast ®lms) and can induce the formation of an interphasial layer close to the interface with the substrate. Considering now the effect of the VA content on the increase of the glass transition at low thickness, an important increase of the Tg is observed for the EVA 14%VA for thickness about 70 mm. For EVA 28%VA and 40%VA, no increase of Tg is seen at the same thickness range. The formation of an interphase of reduced mobility also appears to be favoured in the case of EVA 14%VA. A temperature of about 78C is reached for an EVA 14%VA, 70 mm thick and for a thickness close to only 1 mm for both EVA 28%VA and 40%VA. The greater increase of Tg for EVA 14%VA could be explained by the fact that this EVA is the most crystalline. The mobility of the amorphous chains will be consequently more sensitive to the least modi®cation of the crystallinity organisation. Furthermore, it is interesting to notice that an identical value of Tg, close to 7±108C is reached at low thickness for all the EVA. Indeed, despite the fact that no layer thinner than 30 mm can be obtained for EVA 14%VA, the glass-transition temperature seems to stabilise at a maximum close to 88C, no difference in Tg being observed between the 70 and 30 mm thick ®lms. These results indicate that the minimum mobility of the interphase chains is reached and is independent of the VA content of the EVA. But nevertheless, some additional studies on thinner ®lms are necessary to con®rm the stabilisation of the Tg value for lower thicknesses.
55
4. Conclusion The study of the glass-transition temperature of EVA ®lms deposited on a glass substrate by dielectric spectroscopy (DETA) leads to the following conclusions: In the experimental conditions used in this study, the observed tan peak corresponds to the glass-transition relaxation. The temperature corresponding to the maximum of the peak is related to the glass-transition temperature, Tg. The results obtained by DETA show that the glasstransition temperature is significantly shifted towards higher values when the polymer thickness is decreased, whatever is the VA content. These results indicate that the segmental mobility of the polymer chains is reduced at lower thicknesses, that corresponds to a mobility reduction in the vicinity of the interface with the substrate. These phenomena characterise the existence of a constrained interphase, localised close to the substrate and where the polymer properties differ from the bulk ones. Influence of a confinement effect due to mechanical strains can be rejected. No major effects of the compression process under heating press (used for the EVA 14%VA) or the solvent casting way (used for EVA 28%VA and 40%VA) have been noticed: both processes lead to films presenting an increase of Tg at lower thickness. Finally, dielectric spectroscopy appears to be a easier, sensitive and fruitful technique for studying relaxation phenomena in thin polymer films deposited on a substrate. References [1] K. Varadarajan, J. Coat. Technol. 55704 (1983) 95. [2] N.E. Hill, W.E. Vaughan, A.H. Price, M. Davies, Dielectric Properties and Molecular Behaviour, Van Nostrand Reinhold, London, 1969. [3] A.K. Jonscher, Dielectric Relaxation in Solids, Chelsea Dielectric Press, London, 1983. [4] N.G. Mc Crum, B.E. Read, G. Williams, in: Anelastic and Dielectric Effects in Polymeric Solids, Dover, New York, 1991. [5] O. Pelissou, R. Diaz Calleja, L. Gargallo, D. Radic, Polymer 3516 (1994) 3449. [6] K. Deutsch, E.A.W. Hoff, W. Reddish, J. Polym. Sci. 13 (1954) 565. [7] K.L. Ngai, S. Mashimo, G. Fytas, Macromolecules 21 (1988) 3030. [8] J. Cormenero, A. Alegria, P.G. Santangelo, K.L. Ngai, C.M. Roland, Macromolecules 27 (1994) 407. [9] J.V. Standish, H. LeidheiserJr., J. Coat. Technol. 53678 (1981) 53. [10] R.G. Dillingham, F.J. Boerio, J. Adhesion 24 (1987) 315. [11] J.S. Crompton, J. Mater. Sci. 24 (1989) 1575. [12] H. Schonhorn, J. Polym. Sci. 5 (1967) 919. [13] M. Nardin, E.M. Asloun, F. Muller, J. Schultz, Polym. Adv. Technol. 2 (1991) 161. [14] C.A. Kumins, C.J. Knauss, R.J. Ruch, Progress Org. Coatings 28 (1996) 17.
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[15] J. Brandrup, E.H. Immergut, Polymer Handbook, 2nd ed., Wiley, New York, 1975. [16] R.J. Koopmans, R. Van Der Linden, E.F. Vansant, J. Adhesion 11 (1980) 191. [17] L.E. Nielsen, J. Polym. Sci. 13 (1960) 357. [18] D.E. Buerger, R.H. Boyd, Macromolecules 22 (1989) 2699. [19] H. Braunisch, Acustica 22 (1969-1970) 136.
[20] S. Bistac, D. Cheret, M.F. Vallat, J. Schultz, J. Applied Polym. Sci. 65 (1997) 347. [21] H. Bourhala, J. Chauchard, J. Lenoir, H. Romand, Die Angewandte Makromolecular Chemie 1782941 (1990) 47. [22] O.N. Tretinnikov, R.G. Zhabankov, J. Mater. Sci. Lett. 10 (1991) 1032.