thermal and electrical properties of some epoxy based composites

Thermal and Electrical Properties of Some Epoxy Based Composites

GYORGY BANHEGYI

Research Institute for the Plastics Industry Budapest, Hungary

and

PAL SZAPLONCZAY and GABOR FROJIMOVICS Research Institute for the Electrical Industry

Budapest, Hungary

and

FRANK E. KARASZ Polymer Science and Engineering Department

University of Massachusetts Amherst, Massachusetts 01 003

Thermally stimulated creep compliance, differential scanning calorimetric behavior, thermal degradation, AC dielectric permittivity and loss (between 120 Hz and 1 00 kHz) and thermally stimulated polarization and depolarization currents were studied in a cycloaliphatic epoxy resin (B), in a conventional bispheno1-A- diglycidyl ester type epoxy resin (C) and in composites consisting of: resin B/ wollastonite (B/W), resin B/quartz (B/Q) and resin C/wollastonite (C /W) . The filler content was 60 wt%. Resin B exhibited higher Tg and lower rubbery deformability than resin C due to its more compact structure. Fillers reduced the rubbery deformation and thermal expansion and shifted the transition temperatures by a few degrees. The shift depended on the method used. Composites B/W and C/W exhibited higher' thermal stability than the corresponding pure resins, while sample B/Q was less stable than resin B. Resins B and C exhibited a low temperature /3 transition (in the case of resin B a doublet) and a high temperature a or glass transition. AC dielectric losses were fairly similar in samples B and B/ W, while the high temperature loss of sample B/Q was determined by a space charge process probably due to the matrix/filler interface. In samples C and C/W the QI transition is visible but it is superposed on a strong space charge process due to the resin/electrode interface. Thermally stimulated currents show a behavior qualitatively in agreement with the AC results but the very low effective frequency and the nonlinear field strength dependence of the space charge processes cause some minor differences.

INTRODUCTION

he importance of epoxy resins in the electrical T and electronic industries is well known. They serve as encapsulating and potting resins, coatings, adhesives, printed circuit board materials, etc. (see e.g. Refs. 1-3). They form a versatile family of plastics due to the variability of the epoxy containing components and hardeners, or curing catalysts. Their chemistry is briefly reviewed in Refs. 1-4. To reduce the shrinkage and warpage on curing, to im- prove heat distortion and heat dissipation, and to match the heat expansion of the encapsulating ma-

terial to that of the encapsulated device, mineral fillers are usually added to the epoxy, occasionally at a high loading level.

In this article we discuss some thermal (thermomechanical, calorimetric and heat degradation) and electrical (AC dielectric permittivity and loss, DC thermally stimulated polarization and depolarization) properties of some highly filled epoxy composites and compare them to the corresponding properties of the pure matrix materials.

One of the main purposes of this study was to explore the properties of a newly developed wollastonite filler and to compare its properties to a conven-

POLYMER COMPOSITES, JUNE 7990, Yo/. 77, No. 3 133

G. Banhegyi. P. Szaplonczay, G. Frojimovics, and F. E. Karasz

tional silica flour filler widely used in the electrical industry.

EXPERIMENTAL

Samples

Two different resins were used in this study. One was the product of Bayer (hereafter denoted resin B) obtained by mixing 100 wt Lekutherm XlOO cycloaliphatic diglycidyl ester (epoxy equivalent about 1 70- 180) with 100 wt Lekutherm Harter M (methyl hex- ahydrophthalic acid anhydride, MHHPA); the accel- erator was DMP-30 (2,4,6-tris(N,N-dimethyla- mino)phenol), 0.4 wt.

The other resin was the product of Ciba Geigy (hereafter denoted resin C) consisting of 100 wt CY225 modified epoxy-Bisphenol-A and 80 wt HY925 anhydride type crosslinking agent.

Resin B was filled with 60 wt% W12 EST type silica flour (crystalline, 90% of the particles between 2 and 64 pm in diameter, product of Quartzwerke, Frechen, FRG) (hereafter denoted composite B/Q) and also with 60 wt% wollastonite (shape factor L/D about 10-30, diameter of 64% of the particles between 2 and 16 pm, product of Quartzwerke, Frechen, FRG) (hereafter denoted composite B/W). Resin C was filled with 60 wt% wollastonite of the same type used in composite B/W (hereafter denoted composite C/W) . Com- posite B/Q also contained about 0.4 wt% Fez03 (Bay- ferrox Eisenoxid Rot 130) which is generally used as a coloring agent in epoxy formulations. Quartz and wollastonite fillers were surface treated with epoxy- silanes. The fillers were dried at 50°C in a 5-10 torr vacuum (about low3 MPa) to remove adsorbed water before mixing with the resin. However, the composite may itself absorb water under ambient conditions.

The components were mixed at room temperature, the tool temperature in the press was 14OOC. applied pressure was 0.2-0.3 MPa, molding time was 10- 15 min. The samples were postcured for 12 h at 120°C. Pure resin samples were obtained under similar thermal conditions, but between glass plates without pressure.

Measurements

The thermomechanical properties of the samples were studied using a microcomputer controlled device called MULTIRELAX designed by Hedvig (5). The principle of the measurements is as follows (5, 6): four samples are placed into a thermostated copper block, and their deformation is studied under load. The load is applied through a balance arm either in extensional or compressional mode. The load level and deformation is recorded as a function of temperature on heating (usually l"C/min). The samples are usually studied in pairs: one loaded while the other one is unloaded (in fact a very small load is applied to obtain contact with the sample and to overcome some friction in the balance); thus, the true deformation of the sample can be obtained as the difference between the loaded and unloaded channels. The deformation of the unloaded channel is the sum of

thermal expansion and the recovery of "frozen-in" deformations which are due partly to shrinkage upon crosslinking, partly to physical aging and partly to unavoidable local orientations occurring during sample preparation. The thermally stimulated creep compliance of the system is calculated as a quotient of the true deformation and the load.

DSC and thermogravimetric (TGA) curves of the systems were studied using the Perkin-Elmer 7 Ther- mal Analysis System. In both cases the heating rate was 20°C/min. TGA was performed in air.

The AC dielectric properties of the samples were measured using a Hewlett-Packard 4274A type RLC digital bridge in the 100 Hz- 100 kHz frequency range. Samples 2-3 mm thick (in the case of the composites) or about 1 mm thick (in the case of the pure resins) and 55 mm in diameter were equipped with evaporated A1 electrodes (diameter 40 mm) and were placed in a thermosiated block (7). Surface currents were eliminated by an evaporated guard ring. Dielec- triq parameters were calculated using the room temperature geometric capacitance; thermal expansion was not taken into account. Capacitance and tan S were recorded at 100 kHz, 10 kHz, 1 kHz, and 120 Hz. In the last case 120 Hz was chosen to avoid interference with the first overtone with the power frequency (50 Hz), but the 120 Hz loss values were not consistent except above a certain loss level (about lo-' tan 6). Curves were measured upon both heating and cooling (2"C/min) between - 100 and +200"C.

Thermally stimulated DC measurements were performed using a custom made system designed by Hedvig (7) as a part of the MULTIRELAX system. With this system currents between and A can be detected with a maximum time constant of 1 s. The polarizing voltage can be varied between 10 V and 2 kV. In the thermally stimulated polarization (TSP) measurements the samples were polarized at room temperature (5 kV/cm, 30 min) and the current was recorded as a function of temperature on heating up to 150°C. The TSP current is a sum of the ohmic conduction current and the polarization current. At 150°C the samples were further polarized for 30 min, cooled under the effect of the field to -80°C; the external field was then removed and the short circuit current was detected on reheating. The latter is re- ferred to as thermally stimulated depolarization (TSD) current. All heating and cooling rates in these measurements were 2"C/min.

Results

The thermomechanical properties of the pure and filled epoxy systems will be discussed first. Figure 1 shows the deformations of the unloaded channels between room temperature and 190°C. Several ob- servations can be made:

a) both resins B and C in the pure state have some "frozen-in" deformation which relaxes at the glass transition temperature (note the S-like curvature in the thermal expansion curves); b) The thermal expansion coefficient of the pure resin (and conse-

134 POLYMER COMPOSITES, JUNE 1990, Vol. 1 I, No. 3

Thermal a n d Electrical Properties of S o m e Epoxy B a s e d Composi tes

/

Fig. 1 . Thermal expansion of resins B and C and composites B/ W, BIQ, and C/ W.

quently, of the composites) increases on passing the glass transition temperature; c) The addition of fillers decreases the “frozen-in” deformation and the thermal expansion coefficients, especially below Tg. The acicular wollastonite filler (which can be approxi- mated by a prolate ellipsoid) decreases the thermal dilatation coefficient more effectively, than Lhe nearly spherical quartz filler in the case of resin B.

The extensional thermomechanical curves (calculated from the true deformation) of the two pure resins are shown in Fig. 2. The deformability of the sample is compared to the room temperature value. The temperature derivatives are also shown, which helps to identify the transition temperatures. Both curves are typical for thermoset networks; at the glass transition the compliance increases considerably, but at higher temperatures it reaches a rubbery plateau and does not change until the chemical decomposition of the network begins. This did not occur up to 200°C in this case. The glass transition temperature of resin C occurs about 2OoC lower than that of resin B and the rubbery deformability plateau is higher in resin C indicating that its network is looser, provided that the chains between the crosslinks are not too short and the chain flexibilities are comparable.

The thermomechanical curves of the filled samples are shown in Fig. 3. It is apparent that the rubbery deformation above Tg is reduced significantly as compared to the pure resins. In the case of identical fillers (composites B/W and C/W containing 60 wt% wollastonite) the rubbery deformation is higher with resin C than with resin B, similar to the pure resin components. If the resin component is identical (compos-

POLYMER COMPOSITES, JUNE 1990, VOl. 1 1, NO. 3

20 LO 60 80 100 120 1LO 160 180 2CQ Fig. 2. Thermally stimulated creep compliance curves measured from room temperature for resins B and C . The temperature derivatives of the curves are also shown.

1 0.010

I I I I I 1 1 1 I I 20 Lo 60 80 XK) m uo 160 180 200

Fig. 3. Thermally stimulated creep compliance curues and their temperature deriuatiues f o r the composite samples.

135

G. Banhegyi , P . Szaplonczay , G . Frojimouics. and F. E. Karasz

ites B/g and B/W containing 60 wt% quartz and wollastonite, respectively) then the acicular wollastonite filler reduces the rubbery deformability more effectively than the more isotropic quartz filler.

The transition temperatures obtained by the thermomechanical method and DSC are compared in Table 1. The agreement between the two sets of data is acceptable but not quantitative. The addition of fillers clearly shifts the transitions but the two methods show opposing shifts in some cases. Thus, when discussing the "reinforcing effect" of a filler in terms of glass transition temperature shifts, one must be cautious about which method is chosen. The DSC curves did not exhibit significant endotherms on heating above the glass transition temperature which would indicate gross undercuring. However, the possibility that higher temperature postcuring causes further crosslinking is not ruled out [see the dielectric data shown later).

The thermogravimetric curves of the resins and composites are shown in Fig. 4 and the decomposition temperatures [the maxima of the differential thermogravimetric curves) are listed in Table 2. There is practically no weight loss up to 250°C [less than 5%). The thermal decomposition proceeds in two steps. In the case of resin B the second step is not as well defined as in resin C. This two-step char- acter can also be observed in the composite samples but the maxima shift with respect to the pure polymers. The wollastonite filler increases the first decomposition temperature in both cases, while in the quartz filled sample the system begins to decompose at lower temperatures. [This may also be due to the Fe203 additive which is present in small amounts). The second decomposition temperatures shift toward the lower values in all composites.

Figures 5 and 6 show logarithmic plots of the dielectric losses measured in pure resins B and C, respectively. The presence of at least two transitions is clearly observed in both samples. In epoxy resins the glass transition is usually denoted the a transition, the second transition, appearing between 0 and -50°C is the 0 transition, and in some cases a third, so-called y transition, can also be detected at cryogenic temperatures (7-13). The fl transition is broader and of much lower magnitude than the a transition. In sample B and in its composites the p transition appears as a doublet (see Fig. 7). The onset of DC conductivity at high temperatures and low frequencies is more dominant in resin C. in accord-

Table 1. Thermomechanical and DSC Transition Temperatures of the Pure Epoxy Resins and their Composites.

Transition Temperature ("C)

Sample DSC Thermomechanical

B 116 122 112 118 105 124

B/Q

96 98 B/w C c/w 95 114

100

00

60 8

'3 3 4c

I- I

W

20

0

I I

I I I I I I L

300 400 500 6 00

TEMPERATURE ( "C ) Fig. 4. Thermogravimetric curves of the resins and compos i tes .

Table 2. Thermal Decomposition Temperatures (Maxima of the DTG Curves) of the Pure Epoxy Resins and their

Composites. (Measured in air, heating rate 20°C/min).

Sample T, ("C) 1 2 ("C)

B 440 535 395 450 455 495

B/Q

415 585 B/W C ClW 450 535

ance with its lower Tg and looser network structure [see Fig. 2). The non-exponential growth of the 120 Hz loss in sample C at the highest temperatures indicates that it cannot be simply ascribed to conduction.

Figures 8, 9, and 10 show the dielectric losses measured in composites B/W, B/Q, and C/W respectively. Comparing the wollastonite filled B/W sample with pure resin B [Ffgs . 8 and 5, respectively) one can see that at high frequencies the change is significant. The low frequency loss, however, shows con- siderable increase beginning from the pre-glass transition region. Sample B/9 shows markedly different behavior: a dispersive conduction process is superposed on the transitions which dominates the low frequency behavior from about 50°C.

136 POLYMER coMPos/TEs, JUNE 7990, Yo/. 17, No. 3

Thermal and Electrical Properties of S o m e Epoxy Based Compos i t e s

=w m -

-I

- - 2

I I I I

B I - - 10

120 Hz I KHz 10 KHz

- IOOKHz

- - - _. . . .

01

I I I I 0 01 0 200 50 50 100 150

-- _ _ O -

-

..-. />--? - <- - / -+ --._.__-*' - _ -. ---_ , --_

- 1.0

W

- 0 1

0

w m -

-I

- 2

I , I , l I , - 2 1 ' l o o 1 -50 50 100 150 200

TEMPERATURE ( "C )

Fig. 6. Dielectric loss curues of resin C.

' 1 ' 1 ' 1 '

E / W I -

120 Hz I KHz 10 KHz 100 KHz

-.- _ -___-- - -. , - . . - .

0 -

- ,--+A=: -. . _-- ..-.

I I I 1 -

There are also differences between samples C and C/W. The transition becomes broader in the composite and the 120 Hz loss does not tend to level off at the highest temperatures. The 1 kHz transition temperatures of the resins and their composites are summarized in Table 3 together with the corresponding activation energies.

The real parts of the permittivity ( E ' ) for B, B/Q, B/ W and for C, C/W are plotted logarithmically as a function of temperature in Figs. 1 1 and 12, respectively. (The curves are not plotted below 50°C because in the /3 transition range the frequency dependence of c' is very slight, especially in the logarithmic representation). The €' steps due to the a relaxation are clearly present, but in some cases (B/Q, C, C/W) the onset of a new polarization mechanism is also evi- dent. The Cole-Cole representation of the 120 Hz data of these systems (Fig. 13) shows that the increasing loss is accompanied by increasing c' as well, so that the high temperature, low frequency losses cannot be attributed solely to the ohmic conduction. A higher magnification of the Cole-Cole plots constructed from points measured at different frequencies for system B/Q shows (Fig. 14) that while in the (Y relaxation range the points taken at different frequencies can be superimposed to a single curve; in the higher

I&-

Tit) 0 ' ' ' ' ' ~ 1 ' ~ ' ' ' ' ' 1 1 ' '

Fig. 7 . Dielectric loss curves of resin B and its composites [B/Q and BIW) in the p relaxation range. The linear scale shows the peak doubling.

- 50 0 50 100

temperature range where the new mechanism begins to be important this "thermorheological simplicity" is no longer valid. The nature of this high temperature transition will be discussed subsequently.

The dielectric parameters were also measured upon cooling to study possible irreversible changes occurring during the first heating. Figures 15 and 16 show the 1 kHz and 100 kHz loss curves for composites B/ W and C/W, respectively on heating and cooling, as examples. In the case of sample B/W the most important change is the shift of the transition temperature toward higher values, while for sample C/W the de-

POLYMER COMPOSITES, JUNE 1990, Vol. 11, No. 3 137

G. Banhegyi, P. Szaplonczay, G. Frojimovics, and F. E. Karasz

I I I I

8 1 0 I -

10 KHz IOOKHz

. - - - - - - - , - . . - . . - ew O -

0 -

-I - 0.1 -

I I I I 0

TEMPERATURE [ "C )

0 01 50 100 150 200

- 2 -50

Fig. 9. Dielectric loss curves of composite B/Q.

10

10

0 I = 0 1

50 100 150 200 -2 -50

TEMPERATURE ( 'C 1

Fig. 10. Dielectric loss curves of composite C/W.

Table 3. Dielectric Relaxation Peak Temperature at 1 kHz and Corresponding Activation Energies in the Pure Resins and in

the Composites.

0 transition@) (Y transition

HL? WJI H, Sample ("'I mole) T, ("C) (kJ/mole)

B -38 -18 45 52 147 264 B/Q -55 (?) 54 - 134 247 BpN -45 5 45 60 135 238 C -48 52 116 229 C/W -36 41 120 281

crease of the low frequency loss is also significant. In the latter case the low frequency permittivity also decreased considerably on cooling.

Finally, the thermally stimulated M3 properties will be presented. Figure 17 shows the thermally stimulated polarization [TSP) curves for the resins and for the composites in an Arrhenius representation. It must be noted that this curve is not simply the conductivity curve of the material, it is rather the sum of the conductivity and other polarization contribu- tions (7). In the presence of dipolar or space charge mechanisms this difference becomes important. In

I I

_ _ _ _ _ - - - - - -

0.7 I .3

. e / Q

1 . 1 -

- - - - - - - I KHz

0.9 - 10 KHz - . . - , . -

0.7 -

0.5 -

5 0 I00 150 200

TEMPERATURE ( "C Fig. 11. Dielectric dispersion curves of resin B and composites.

I .5

I .O

its

I I I I I

TEMPERATURE ( OC) Fig. 12. Dielectric dispersion curves of resin C and its composites.

the case of resin B [which exhibits lower conductivity than resin C, especially in the high temperature range), the conductivity of the composites is higher than that of the matrix resin. In the case of resin C

138 POLYMER COMfOSlTES, JUN€ 1990, Vol. 11, No. 3


€' ( 120Hz 1 Fig. 13. Cole-Cole representation of the temperature dependent 120 H z data of those samples which exhibit space charge polarization.

r I I I I I I I (

4 5 6 7 8

€' Fig. 14. Cole-Cole representation of the dielectric data taken at dwerent frequencies for sample B/Q in the a relaxation range and above.

the conductivity curves of the matrix resin and of the composite cross each other: at low temperatures the conductivity of the matrix is lower, at higher temperatures the situation is reversed. It is worth mention-

log c"

. . . . . . . 100 kHz

0 'I

-150 -KO -50 0 50 100 I 5 0 200

Fig. 15. 1 kHz and 100 kHz dielectric losses of sample Bj Won heating and cooling.

-1 .

-1% -100 -50 0 50 100 150 200

Fig. 16. 1 kHz and 100 k H z dielectric losses of sample C/ Won heating and cooling.

ing that the near equilibrium DC conductivities measured at room temperature are remarkably low. Hand- books usually quote conductivities in the lo-' '-10-13 ohm-'m-' range (1-3) for epoxies, but these are probably either calculated from AC losses or from current values taken after a short pre-set time (usually a few minutes). Figure I8 shows the thermally stimulated depolarization (TSD) curves of the pure resins and of the composites after a 30 min polarization under a voltage of 1 kV at 150°C. The depolarization curves of the pure resins show the glass transition as a

POLYMER COMPOSITES, JUNE 7990, V d . 11, NO. 3 139

G. Banhegyi, P. Szaplonczay, G.

TEMPERATURE ( O C )

200 I50 100 50 0 1 " " I " " 1 ' 3 ' [ 4 1

0 010 0 / w C c / w

I 1 I -1 18 2 .O 2.5 3.0 3 5

1031 TEMPERATURE ( I /K I Fig. 17. Thermally stimulated polarization ITSP) currents of the resins and of the composites in reduced current density units between room temperature and 150°C in the Arrhenius representation.

It

2 -3 i! -I -4

/ - 5 I I I I I I I

-100 -50 0 50 100 150 200

TEMPERATURE ( "C Fig. 18. Thermally stimulated depolarization (TSD) currents of the resins and of the composites fn reduced current density units after a 30 min. polarization at 150°C. EPoI = 5 kV/cm.

maximum and the onset of a new, high temperature mechanism in addition to the broad, shoulder-like /3 transition which appears in this case below room temperature. The high temperature depolarization current is higher in resin C than in resin B, in ac- cordance with the differences observed in their AC behavior [see e.g. Figs. 5 and 6). The depolarizatlon current level is increased by about an order of magnitude in the P relaxation range of the composites as compared to the pure matrices.

Frojimovics. and F. E. Karasz

DISCUSSION

Cycloaliphatic epoxy resins usually exhibit a more compact structure and higher crosslink density than the more widely used bisphenol-A-diglycidyl ester based compounds which results in a higher glass transition temperature and better high temperature stability (3). These general trends can be observed in the present system as well. Resin B, which is a cycloaliphatic epoxy compound for outdoor applica- tions, exhibits a higher Tgr a smaller rubbery deformability (higher modulus) and a lower high temperature conductivity than resin C, which is a more common bisphenol-A-diglycidyl ester type compound. Heat stability under load can be well estimated by the thermomechanical method, which gives information analogous to but more detailed than the heat distortion temperature test. According to this method, addition of quartz or wollastonite filler to resin B hardly changes the deformation temperature, while in the case of resin C wollastonite increases it slightly (see the right hand column of Table 1 ). The DSC transition temperatures and dielectric peak temperatures (see Tables 1 and 3) indicate a decrease in the glass transition temperature in the composites based on resin B as compared to the matrix compound, and a very slight increase in composite C/W as compared to resin C.

The heat decomposition temperature (see Table 2) is positively affected by the addition of the wollastonite filler, and some decrease in the decomposition temperature is observed in resin B on the addition of quartz.

Dielectric studies reveal more structural details and add valuable complementary information for the electrical engineer. Another observation is the double /3 peak in resin B as compared to resin C. The /3 peak is usually attributed to the -O-CH2-CH(OH)- CH2- unit (see e.g. Refs. 1 1 , 14). Chang and coworkers (12, 13) attributed this relaxation to the

structural unit comparing a wide variety of epoxy resins. I t may be assumed that resin B contains two different units of this type while resin C is structur- ally more homogeneous. The p activation energies listed in Table 3 are comparable to those obtained by Chang, et al. (12). (30-50 kJ/mol) using the TSD technique. In fact, these authors have shown by the initial rise method that the activation energy increases continuously in the cryogenic temperature range from about 10 kJ/mol to 80 kJ/mol. The inter- mediate value quoted above can thus be considered as a kind of average. Shimbo and coworkers (1 5. 16) have thoroughly studied the effect of chemical structure on the /3 relaxation properties of anhydride-cured

140 POLYMER COMPOSITES, JUNE 1990, Yo/. 11, No. 3

Thermal and Electrical Properties of S o m e Epoxy Based Compos i t e s

epoxies using the torsional pendulum technique. They have, also, estimated the activation energy by integrating the temperature dependent mechanical loss peak and obtained values (70-90 kJ/mol) which are higher than ours, close to the upper limit given by the TSD technique (12). The method applied by Shimbo, et al. (1 5), to estimate the activation energies involves an extrapolation of the limiting moduli to the peak temperature which gives an uncertainty to the absolute value. In our opinion activation energies based on peak shifts measured at different frequencies are more reliable.

It is interesting to observe that the presence of fillers influences the transition temperature and activation energy of the @-transition which are due to some local segmental motion in the chain. The filler loading is, however, high and the epoxy-silane coupling agents can give rise to chemical bonds at the filler/resin interface, so even local motions can be affected. It is worth noting that in the case of the @- transition the transition temperature and the activation energy do not always change in the same direction as intuition would suggest.

The (Y transition can be identified with the glass- rubber transition. In resin B, in spite of chemical coupling, filling decreases somewhat both the transition temperature and the activation energy of the a transition. In the case of resin C filling is accompanied by a slight increase of T, and a significant increase of the activation energy. According to the adsorption theory of filler effects (17) the interaction of the filler and the polymer molecules can both increase and decrease the transition temperatures depending on the relative importance of two effects. If the adsorption is very strong, a dense interface appears between the matrix and the filler with higher Tg and lower mobility. If the filler loading is high the interface constitutes the major part of the matrix, the TB of the sample is increased. If, however, the interaction is weaker, another effect becomes important: the adsorption points of the macromolecule be- come fixed, the end-to-end distance increases, the adsorbed layer becomes looser, which leads to lower Tg and activation energy in the interface. Seemingly, this is the case with resin B which has a stiffer structure than resin C.

The other interesting problem is the appearance and the nature of the high temperature relaxation processes above the (Y transition temperature. Their presence is most conspicuous in the 120 Hz permittivity and loss curves. The appearance of a low frequency relaxation phenomenon due to the interfacial- or Maxwell-Wagner polarization process can be expected in heterogeneous systems such as polymeric composites (7). In the same temperature and frequency region, however, another electrical relaxation process can appear, i.e. electrode polarization. This is a complex phenomenon, not well understood in polymers. In semiconductors, space charge injection can significantly change the electrical behavior of the bulk material (1 8). If the charge carriers are ions and the electrodes are partially blocking in nature,

POLYMER COMPOSITES, JUNE 1990, Yo/. 7 7 , No. 3

space charge due to the nontotal neutralization of charge carriers (see e.g. Refs. 19 and 20) can appear. Since little is known about the conduction mechanism in polymers (see e.g. Refs. 21, 22), it is hard to decide between the various mechanisms. The presence of interfacial polarization can be proven in fa- vorable cases using theoretical formulae and numer- ical methods (see e.g. ReJ 23) but this requires the exact knowledge of the electrical properties of the matrix and of the filler, and is complicated by the fact that the electrical properties of thermosets can change considerably on curing and postcuring (7.24- 26). In the present case we use a simple experimental method to clarify the nature of the high temperature- low frequency relaxation process: if the pure matrix resin does not exhibit high temperature dispersion, but its composite does, the process can probably be ascribed to the matrix-filler interface; if, however, the pure resin itself produces significant high temperature polarization, it can probably be attributed to the polymer-electrode interface.

In resin B the 120 Hz dielectric loss does increase at the highest temperatures (see Fig. 5) but it is not accompanied by a significant increase of e’ (cf. the lower part of Fig. 111; thus, it can be attributed to a conduction process. The conduction is presumably ionic, and it is quite possible that at even higher temperatures the electrode polarization process becomes dominant. The slight but noticeable increase of E’ (1 20 Hz) at the highest temperatures suggests this possibility.

Resin C, on the other hand, exhibits a well defined electrode polarization process in the unfilled state. The leveling off of the c“ (120 Hz) curve (Fig. 6) . the monotonic increase of the log t” (120 Hz) curve with a shoulder (see the lower part of Fig. 12) and the shape of the Cole-Cole curve (Fig. 13) clearly show that the high temperature process is not simply conduction, but that strong space charge effects are involved.

If one investigates the dielectric properties of the composites (Figs. 8-13), the following features can be observed: 1) The high frequency, low temperature (unrelaxed) permittivity of the composites is always higher than that of the matrix compound. This is due to the fact that the permittivity of mineral fillers is usually higher than that of the epoxies (t‘ = 3.8 for crystalline quartz (27) and t’ = 6.5-8.6 for wollastonite depending on the firing temperature and crystalline modification (28); 2) The difference between the relaxed and unrelaxed permittivities in the transition range decreases in the composites due to their lower polymer content; 3) The case of composite B/W is qualitatively similar to that of pure B matrix, but the a transition shifts towards higher temperatures and the loss is higher in the /3 relaxation range. The low frequency (1 20 Hz) loss of the composite exceeds that of the matrix compound, but there is no sign of significant interfacial polarization at either the polymer filler or at the polymer/electrode interface: 4) In the case of composite B/Q the higher frequency (100 kHz) loss curves of the composite and that of the

141

G. Banhegyi, P . Szaplonczay, G. Frojimovics, and F. E. Karasz

matrix material differ little indicating " excess" loss even at the highest frequency, since the polymer content of the composite is significantly lower. At lower frequencies, however, the differences are significant. The 01 transition process is almost masked by the presence of a broad dissipative process beginning in the @ transition range (below 50OC). The high temperature, low frequency losses of the composites exceed considerably those of the matrix compound and e ' also increases sharply above 170OC. Since no significant electrode polarization process was observed in the matrix compound below 200°C, this change can be attributed to the presence of filler. At first glance this is unexpected because dry crystalline quartz shows very low conductivity (about 1 0-13 ohm-' m-', that of imperfectly dried quartz varies between 10-7-10-13 ohm-' m-', see Ref. 25) and strong interfacial polarization can be expected if the conductivity of the inclusions exceeds that of the matrix by several orders of magnitude (see e.g. Refs. 7 or 23). On the other hand, finely ground silica flour tends to pick up chemisorbed water removable only by extended baking and after which some surface Si- OH groups can remain, causing enhanced surface conductivity (see e.g. Ref. 29 on the dielectric effect of sorbed water). A conducting surface layer can lead to the same interfacial polarization phenomenon as if the whole particle were conductive (30). The possibility cannot be excluded that the desorbing water or the dissociated protons of the silanol groups increase the charge carrier concentration in these composites, leading to an electrode polarization process. The Fe,Os coloring agent might also have some interfacial effect since transition metal oxides are usually semiconductive, but the volume fraction of this component is so low that it alone cannot cause the observed space charge effect; 5) Composite C/W exhibits a space charge polarization process similar to that of the matrix compound, with two notable differences: both the a and 6 transitions shift toward higher temperatures; the low frequency loss curves increase more sharply above the a transition temperature in the composite. This latter feature can be seen in Fig. 13, which compares the Cole-Cole curves of samples exhibiting a space charge relaxation. The shape of the curve of sample C/W suggests that in this case the matrix-filler and the matrix-electrode polarization processes are superposed onto each other and on the conductivity.

The non-dipolar nature of the high temperature relaxation processes is also proven by their thermo- rheologically complex behavior (see e.g. Fig. 14 for composite B/Q).

The heating-cooling cycle has a double eGect on the electrical properties (see Figs. 15 and 16): the a transition temperatures increase which might be due to a post-curing effect (the samples have been postcured at 12OOC -just prior to the measurement) and/ or to the removal of a small amount of water which acts as a plasticizer and affects the polymer-filler interface (see Refs. 14 and 31-36); if the composite

exhibits @pace charge polarization, its magnitude decreases (both t' and c") due to post-curing, thus, reducing the mobility of all charge carriers in the matrix.

The present findings on the thermally stimulated DC behavior are in agreement with the AC properties and with earlier studies on epoxy-mica composites (37). In systems based on resin B the tendencies are simple and clear: the thermally stimulated polarization currents of samples B/Q and B/W exceed those of sample B in the whole temperature range: the difference is more important in the low temperature range. The TSP curve of sample B/Q is higher than that of B/W, similar to the AC loss curves. In the depolarization regime pure resin B shows a very broad B transition, a sharper 01 transition and the onset of a new, space charge mechanism. In samples B/W and B/Q the TSD current level is considerably increased (by about one order of magnitude at low temperatures and more at high temperatures), the two curves are very close to each other except in the highest temperature range where the TSD current of composite B/Q is higher. This similarity of the B/Q and B/W TSD curves seems to be in contrast to the AC loss behavior where sample B/Q was markedly different from samples B/W and B. One has to take into account, however, the big difference in the field strengths used in these two methods: the measuring voltage in the AC measurements is 1 V, while in the DC measurements it is in the range of 1 kV. Nonlin- earity of the various space charge processes is probably responsible for the observed differences. In- creased TSD currents in the /3 relaxation range, where neither space charge, nor the classical interfacial polarization mechanism can be active, are probably due to the increased defect concentration (38) introduced by filling. In composites B/W and B/ 9 the high temperature space charge process totally masks the a transition of the matrix compound; only a single broad high temperature process can be seeii. In samples C and C/W the situation is somewhat different. In the low temperature range both TSP and TSD curves of composite C/W exceed those of resin C. At high temperatures the thermally stimulated conductivity of resin C exceeds that of composite C/ W, behavior different from that observed in the AC measurements. This is again probably due to the different field strength dependence of the various space charge processes (e.g. charge injection, electrode polarization and interfacial polarization). The 01 depolarization peak of resin C is much higher than that of resin B, because it is superposed on a very intensive space charge peak which is much less pro- nounced in resin B. The a transition of resin C can still be seen as a shoulder on the space charge peak in composite C/W. The fact that the depolarization curves of C and C/W practically coincide at high temperatures show that the excess thermally stimulated conductivity of C as compared to C/W is due mainly not to additional polarization but to more intensive charge transport.

142 POLYMER COMPOSITES, JUNE 1990, Vol. 11, No. 3


SUMMARY AND CONCLUSIONS

Electrical characterization methods, in combination with other thermoanalytical methods, proved to be useful in comparing the properties of epoxy based composites with those of the matrix compounds. The most important findings can be summarized as follows: 1) The addition of fillers decreases the thermal expansion, the high temperature deformability and the “frozen-in” deformation in the composites. There are slight, but detectable changes in the glass transition temperatures, but the different methods do not always give the same shifts; 2) The AC dielectric loss in the ,8 relaxation range shows little change in the composites as compared to the resin compounds. At high temperatures above Tg new polarization mechanisms appear in samples B/Q, C, and C/W, which are probably due to the matrix/filler interface in the first case and to the matrix/electrode interface in the latter two cases. In the last case the admixture of matrix/filler and electrode polarization mechanisms is possible. These processes are frequently nonlinear in nature and are due to mobile charge carriers (here probably ions). Post-curing effects can be observed on cooling: Tg increases, the magnitude of the space charge processes decreases: 3) The thermally stimulated current experiments usually corroborate the results obtained by AC dielectric methods. The differences appearing in some cases are probably due to the different field strength dependence of the various mechanisms involved. Especially notable is the enhanced depolarization current level at the lowest temperatures, which shows the importance of “defect dipole” mechanisms (7, 38) in TSD measurements.

This study proves the necessity of the combination of various electrical and other methods in elucidating the complex behavior of thermoset based composites.

ACKNOWLEDGMENTS

One of us (F. E. K.) would like to acknowledge a grant from AFOSR. %88-001.

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thermal and electrical properties of some epoxy based composites

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