Melt Rheology of Two Engineering Thermoplastics - Highly Filled

Melt Rheology of Two Engineering

Thermoplastics: Poly(ether Imide) and

Poly(2,6-Dimethyl-1,4-phenylene Ether)

DILHAN M. KALYON, DONG-WOO YU, and JEONG S. YU, Department of Chemistry and Chemical Engineering, Stevens

Institute of Technology, Castle Point, Hoboken, New Jersey 07030

Synopsis

Material functions of two engineering plastics [a poly(phenylene ether) and a polycether imide)] were characterized, including the shear viscosity, first normal stress coeffkient, storage and loss moduli, growth and relaxation of shear stress, and first normal stress coeffkient and relaxation moduli. The oscillatory shear and relaxation moduli data were employed to determine the temperature-dependent parameters of Wagner model. Various material functions, which were determined on the basis of this model in conjunction with the fitted parameters agreed reasonably well with the experimental results. The reported data and parameters should facilitate a better understanding of the processability charac- teristics of these two engineering plastics.

INTRODUCTION

Engineering thermoplastics are high-performance resins which can be shaped into products exhibiting excellent ultimate properties, including high rigidity, tensile and impact strength. The processing windows of engineering plastics may be narrow, and detailed rheological characterization can help to define the optimum processing conditions. A number of studies have been reportedie6 on the rheology of engineering plastics and their blends, with three focusing on poly(phenylene ether) resins and their miscible blends with polystyrene.1-3

In the following, we will report our experimental findings on the melt rheology of poly(2,6-dimethyl-1,4-phenylene ether) and poly(ether imide). For brevity, poly(2,6-dimethyl-1,4-phenylene ether) will be referred to as poly(phenylene ether). The side-by-

1988 by The Society of Rheology, Inc. Published by John Wiley & Sons, Inc. Journal of Rheology, 32(B), ‘789-311 (1988) CCC 0148~6055/88/080789-23$04.00

790 KALYON, YU, AND YU

side evaluation of these two engineering plastics was prompted by the observation that the residual stresses found in articles processed from these two resins differed significantly. The rheological characterization is thus a part of a larger study, which encompasses the moldability behavior, microstructure, and ultimate properties of articles processed from these two engineering thermoplastic resins.7’8

MATERIALS

Poly(phenylene ether) is a linear thermoplastic resin with a free phenolic hydroxyl group at the head of each polymer chain and is produced by the oxidative coupling of 2,6-dimethyl phe- nol.gT1o This polymer:

f-Q-f / \ cH3g - CH3 n

has a glass transition temperature of 205°C. It exhibits a very small fraction crystallinity upon being cooled from the melt. Its melting temperature is between 262°C and 267°C. The poly(phenylene ether), which was employed in our study, had a number-average molecular weight of 20,000, a weight-average molecular weight of 48,000, and an intrinsic viscosity of 0.46 dL/g. These properties were determined in chloroform at 25°C and were reported by the manufacturer.

Poly(phenylene ether) resins are thermally stable in the ab- sence of oxygen at temperatures below 300°C. At temperatures higher than 3OO”C, some gel formation occurs and starting at 4OO”C, an exothermic reaction takes place.‘Op” At temperatures above 250°C in the presence of oxygen, poly(phenylene ether) undergoes a rapid oxidation reaction with the formation of gels and colored byproducts.” These thermal and oxidative phe- nomena were considered in sample preparation and rheological characterization of poly(phenylene ether). For cone and plate experiments, the samples were molded directly into a cone- plate shape in a vacuum compression molder at 290°C. For parallel-plate experiments, discs with a thickness of 1.2 mm were vacuum molded. Gel permeation chromatography (GPC)

RHEOLOGY OF ENGINEERING PLASTICS 791

and gel content analysis with a Soxhlet (according to 494) were carried out on molded specimens to verify were not degraded.

ASTM D- that they

The poly(ether imide)?

is a high temperature engineering thermoplastic with a glass transition temperature of 215°C. We have employed an unmodi- fied grade of poly(ether imide) manufactured by General Elec- tric. It has an intrinsic viscosity of 0.47 dL/g and number- and weight-average molecular weights of 12,000 and 30,000, respectively. The polymer was dried at 60°C for 24 h before rheological characterization to reduce its moisture content to below 0.05%. The sample preparation techniques were similar to those of poly(phenylene ether).

Rheological Characterization

The resins were characterized in terms of their linear viscoelastic properties, storage (G’) and loss (G”) moduli employing a Rheometrics, Model 800 Mechanical Spectrometer in conjunction with 25-mm diameter parallel-plate fixtures. The experiments were carried out in the 250-340°C range. Poly- (phenylene ether) is stable under the employed temperatures, and in an inert gas environment, as shown typically in Fig- ure 1. The small amplitude oscillatory shear experiments cov- ered 0.01 to 100 rps frequency range. The magnitude of the complex viscosity, Iv*1 was determined from:

Iq*(w)l = ((G’/w)~ + (G”/w)‘)‘.~ (1)

The Rheometrics mechanical spectrometer was again used for steady shear, stress growth, and stress relaxation experiments, in conjunction with 25-mm cone and plate fixtures. Further- more, a series of stress relaxation after sudden shearing displacement experiments were carried out in the 5-2000% strain range to determine the relaxation moduli, G(t, yo):

Gk~o) = c,(t)lyo (2a)


2’ 0.0 3.0 6.0 9.0 120 1 5 .O

TIMJ3, MIN.

Fig. 1. Storage and loss moduli of poly(phenylene ether) at 290°C as a function of time.

where r,,,(t) is the shear stress and y0 is the imposed shear strain. In these experiments parallel-plate fixtures were employed. The parallel plate data were corrected by employing:13-l5

W,YR) = G,(t,yR) 1 a ln G,(~,Y~)

1 + z a In YR 1 (2bl

where the apparent relaxation modulus G, (t, yR 1 is given as:

G,(~,YR) = 2m YR 1

di 3~~ and strain at radius R:

OR YR = z

(24

CM)

where 9 is the angular displacement, R and H are the plate radius and the gap setting between the two plates, respectively, and Y is the torque. The rise times involved in the imposition of the strain, y0 were not negligible in our experiments and varied between 0.1 to 0.8 s. The relaxation modulus values were corrected for the rise time on the basis of the procedure suggested by Laun.‘”


The shear viscosity of poly(ether imide) at the higher shear rate range of 10 to 1000 s-l was characterized employing an In- stron Universal Tester, floor model with an Instron Capillary Rheometer, employing a series of capillaries. True shear viscosity values were obtained from Bagley and Rabinowitsch correc- tions 17,18 It was not possible to carry out the capillary rheometry experiments with poly(phenylene ether), because of its affinity to oxidation.

Constitutive Equation

As outlined above, the rheological data could only be collected at relatively low temperatures and deformation rates, while these resins are commonly processed at higher temperatures and deformation rates. To gain insight into the processing behavior of these engineering thermoplastics the experimental data were employed to determine the material-dependent parameters of the following integral-type, network-based equation fitting the general form:”

7 = I = ~M,w,,~2)C;1 + ~,W,,~,)C,lds (3) 0

where C;’ and C, are the Finger and Cauchy tensors, respectively, and M, and M2 are the memory functions which are given as functions of the first, I,, and second, I, invariants of the Finger tensor and the elapsed time s. Here we have followed Wagner’s postulate2’ by setting M, to zero and assuming that the memory function, M, can be expressed as a product of two functions:

Ws, Z,J,> = Mo(sMAJ2) (4)

where h(Z,,Z,) I 1 is the temperature-independent damping function which tends to unity for small deformations and MO(s) is the Lodge’s rubberlike-liquid memory function:21

MO(s) = C (Go,/&) q-4-0 - t’)/&) (5)

where G,i and Xi are the relaxation strength and time, respectively. We have employed a double exponential dependence of the damping function to strain;” namely,

WY) = f exp(-w) + (1 - 17 w-+-~2y) (6)


where f, n,, and n2 are material parameters. Various material functions predicted on the basis of this constitutive equation are reproduced here for convenience in Table I. For strain recovery

TABLE I Various Shear Material Functions According to Wagner Model” in

Conjunction with Eq. (6)

Small Amplitude Oscillatory Shear Flow

(11)

G”(w) = 7 s (12)

Steady-Shear Flow

(13)

(14)

Start-up Flow

[I - exp(-t2,J(l + t,,, - y;,)] (16)

where tl,, 1 + n,jh, =-t

A,

1 + ++A, t2,c =-t

AL

Cessation of Steady-Shear Flow

s-(4?) = c IL f&z A, (1 - fE,,A, (1 + n,jh,)2 + (1 + nzjAJ2 1 exp(-t’hJ

2fG, A,2 + 20 - fG,h,2 (1 + n~th,)~ (1 + n*%Q I

ev-W

(17)

(18)

“Ref. 20.


calculations in shear and extensional flows, this model is modi- fied by the incorporation of a functional operator,23 which is given by the minimum value attained by the damping function in the time interval (t, t’).

RESULTS AND DISCUSSION

The storage and loss moduli of the poly(ether imide) and poly(phenylene ether) are shown in Figures l-3. The data were collected at 290, 310, 330, and 340°C for poly(ether imide) and 270, 290, and 310°C for poly(phenylene ether). To ensure that the samples were stable under our experimental conditions, the oscillatory shear data were collected also as a function of time. For example, the time scan presented in Figure 1 indicates that poly(phenylene ether) was stable at 290°C for at least 12 min upon loading. The typical 95% confidence intervals, which were determined according to Student’s t-distribution are also shown in Figure 4. Assuming that these two engineering thermoplastics are thermorheologically simple fluids, the varia- tion in temperature basically corresponds to a shift in time

0 0 340 t - KIN.11 1112

lo-” lo-’ ld 1G lb? a,w rps

Fig. 2. Storage and loss moduli of polycether imide) collected at 290-340°C and their best fit.

796 KALYON,YU,ANDYU

270t Ed A 290% 60 31ot - EPN.11 &I2

Fig. 3. Storage and loss moduli of poly(phenylene ether) collected at 270- 310°C and their best fit.

Fig. 4. Typical 95% confidence intervals of storage and loss modulus values determined according to Student’s t-distribution of poly(phenylene ether) at 290°C.


scale. Accordingly, all relaxation times change with temperature proportional to the shift factor, ar.24

A,(T) = aTk(To) (7)

The relaxation strengths remain constant and the shift factor is determined from the temperature dependence of magnitude of complex viscosity values determined at low frequency:

ad’) = ev(WR)O/T - l/T,)) (8)

where E, is the activation energy and R is the gas constant. The activation energies for poly(ether imide) and poly(phenylene ether) were determined as 43.8 kcal/g mol and 42.5 kcal/g mol, respectively, on the basis of Eq. (8). A reference temperature, T,, of 290°C was selected for both resins. The discrete relaxation spectra for the two resins were determined from the dynamic data employing a pattern search method which minimizes the objective function, F defined as:

F = $ [((Gi’,exp - G,fi,W:,xp)2 + Wi',,,, - %i,WI:,x,)21 i=1

(9)

where N is the number of data points, (G~,,.,,G&,,) available from the dynamic experiments and G,!fi, and Gtfit denote the best fit values on the basis of Eqs. (11) and (12). The best fit of the storage and loss moduli are shown in Figures 2 and 3 for the discrete spectra given in Table II. The relaxation strengths G,, rep- resent the contribution to rigidity associated with relaxation times which lie in the interval lnh and Inh + dlnh.

TABLE II Relaxation Time A, and Strength G,, at 7’ = 290°C

G,.

AL Poly(phenylene ether) Polycether imide)

1.0 E - 04 2.316 E05 7.838 E05 1.0 E - 03 1.210 E05 3.340 E05 1.0 E - 02 8.025 E04 3.130 E05 1.0 E - 01 3.008 E04 1.599 E05 1.0 E + 00 3.804 E03 2.410 E04 1.0 E + 01 4.968 E02 9.031 E02


The results of the step-strain experiments are shown in Figures 5 and 6. In these experiments the broadest 95% conti- dence intervals were observed for the modulus values determined at small strains and relatively long times (in the 7-10 s range). In this range, the 95% confidence intervals went up to a maximum of -+30% around the mean values reported in Fig- ures 5 and 6. This observed scatter should be related to the relatively small torques measured in this range. At higher strains, the typical 95% confidence intervals were within +8-15% of the reported mean values. In the linear viscoelastic region, the shear relaxation modulus values can be related to the relaxation spectrum through:

G(t) = C Go, exP(-tlh) 1

(10)

10

1C

B -- 1c k

s a

10

l(

I I I I

fyELl- fyELl- T =290 “C T =290 “C

A A

‘6

A 0.8 A 1. A 2.5

. A4. a 5.5

a 8. 4 13.

tk 20.

1 10 time,s

2

Fig. 5. Shear relaxation modulus, G(t, y) of poly(ether imide) at 290°C.


PPE

T 2290 “C

x 0 05

0 1. c, 1.5

. 2.5 0 3.5 o 8. (I 13. 4320.

,2 -, 1 10

time. 5

Fig. 6. Shear relaxation modulus G(t, y) of poly(phenylene ether) at 290°C.

The relaxation moduli determined on the basis of Eq. (10) are indicated by the solid curves in Figures 5 and 6. The relaxation data pertaining to the linear viscoelastic range are further shown in Figures 7 and 8 for poly(ether imide) and poly- (phenylene ether), respectively. The bars in Figures 7 and 8 per- tain to 95% confidence intervals at a strain of 0.8 for poly(ether imide) and at 0.5 for poly(phenylene ether), respectively. The relaxation modulus values, which were determined according to Eq. (10) in the linear viscoelastic range, agree well with those determined employing step strain experiments. The use of the discrete rather than the continuous relaxation spectrum is re- sponsible for the weak oscillations in the predicted values of G(t).

800 KALYUN,~U,ANIJ YU

T= 290 c

Fig. 7. Typical confidence intervals of the relaxation moduli in the linear viscoelastic range for polycether imide) and best fit according to Eq. (10).

“j - Eqn.10 += 0.1

1 q 6 y= y= 0.5 0.7 0.8 y=

3 - EqnlO 0 y= 0.1 0 y= 0.3

YZ “,5

TIME , S

Fig. 8. Typical confidence intervals of the relaxation moduli in the linear viscoelastic range for poly(phenylene ether) and best fit according to Eq. (10).


The relaxation moduli of both resins were determined up to a strain of 20, where the behavior is highly nonlinear. The relaxation moduli were separated into a time-dependent modulus, G(t) determined in the linear viscoelastic range, and a strain- dependent damping function, h(y), on the basis of Eq. (4). The damping function describes the destructive effect of the deformation in the entanglement density of the melt. The damping function is given by the vertical shift required to superimpose the curves of relaxation modulus at various strains on the reference curve representing the linear viscoelastic region. However, especially for poly(phenylene ether) the curves of nonlinear relaxation modulus as a function of time for different strain levels are not parallel for all times, as required by the time- strain separability. The factorization of the nonlinear relaxation modulus requires that the slopes of dG(t, y)/dt be constant over the strain range.

This is further elucidated in Figures 9 and 10, where the best fit of the damping function values collected at 2.3 and 7.3 s are compared for both resins. Mean values of the damping function values were also determined based on the best vertical shift of the nonlinear relaxation modulus curves over the entire experi-

Fig. 9. Strain and time dependence of the damping function values of poly(ether imide) and best fit according to Eq. (6).


-’ 0.0 5.0 10.0 15.0 20.0 25.0

SI’R~N (Y) Fig. LO. Strain and time dependence of the damping function values of

poly(phenylene ether) and best fit according to Eq. (6).

mental time range. The time dependence of the damping function is stronger for poly(phenylene ether) in comparison to poly(ether imide). The predicted material functions are obvi- ously affected by the selection of the short or longer time data for representing the damping function behavior. The results are significantly affected by this selection in the high shear rate range and especially for poly(phenylene ether). We have arbi- trarily selected to employ the mean damping function values. The three parameters of Eq. (6), which describe the strain dependence of the mean damping function values in shear flows, are reported in Table III for the two resins.

Figures 11 and 12 show the typical comparison of the experimental shear stress growth behavior versus the predictions of

f nl n?

TABLE III Parameters in Double Exponential-Type Damping Function

Poly(phenylene ether) Polycether imide)

0.1152 0.1791 0.0149 0.0662 0.5893 0.3229

RHEOLOGY OF ENGINEERING PLASTICS

Fig. 11. Shear stress growth for polycether imide).

803

Fig. 12. Shear stress growth for poly(phenylene ether).


the Wagner model. The bars denote 95% confidence intervals of the data collected. The agreement is not good at short times for either polycether imide) and poly(phenylene ether). The predicted shear viscosity values of poly(ether imide) are compared to the experimental values determined with cone and plate and capillary flows at 340°C in Figure 13. The experimental magnitude of the complex viscosity values are also reported. According to the empirical Cox-Merz relationship,25 the magnitude of the complex viscosity approximates the shear viscosity at corre- sponding values of shear rate and frequency.

Furthermore, the comparisons of experimental and predicted first normal stress values and stress relaxation upon the cessation of steady shear are reported in Figures 14 and 15 and Fig- ures 16 and 17, respectively. Overall, the agreement between the experimental and predicted steady material functions is good. However, for the stress relaxation of poly(phenylene ether) upon steady shear at 0.01 s-’ at 290°C shown in Figure 17, the predicted values fall above the 95% confidence intervals of our experimental findings and thus suggest a slower relaxation of the shearing stress than observed experimentally.

Fig. 13. Experimental and predicted shear viscosity of poly(ether imide).

RHEOLOGY OF ENGINEERING PLASTICS RHEOLOGY OF ENGINEERING PLASTICS 805

Fig. 14. Experimental and predicted first normal stress difference of poly(ether imide).

Fig. 15. Experimental and predicted first normal stress difference of poly(phenylene ether).

806 KALYON, YU, AND YU KALYON, YU, AND YU

_ 0.01 TJ -

-9 O.l ' *I 10-l

' ' ' ' "I lo” 1 f

Fig. 16. Shear stress relaxation upon the cessation of steady shear of poly(ether imide).

Fig. 17. Shear stress relaxation upon the cessation of steady shear of poly(phenylene ether).


The effects of the observed time dependence of the damping function values on the predictions were also considered. This was done by the comparison of the predictions, which were based on the use of the two different sets of damping function values pertaining to 2.3 and 7.3 s. Typically, the values of the predicted shear viscosity values differ by 0.1% and O.O2%“at the shear rate of 0.001 s-l, and by 3% and 10% at 100 s-l for poly(phenylene ether) and poly(ether imide), respectively. The predicted first normal stress difference values differ by about 3% and l.% at 0.01 s-l and by 12 and 14% at 1 s-l for poly(phenylene ether) and poly(ether imide), respectively. In the shear stress growth the differences are about l-4% for poly(phenylene ether) and 0.2- 1% for poly(ether imide). Finally, for shear stress relaxation upon the cessation of steady shear the predicted values differ by about l-6% for poly(ether imide) and 3-13% for poly(phenylene ether) depending on which set of damping function values are used. The differences increase with increasing shear rate at the wall values especially in the greater than 10 s-l range.

The rheological behavior of the two resins can be further employed to gain insight into the processability behavior of the two resins under their pertinent processing conditions. One aspect, which was of concern to us, and which prompted the study, was the relative relaxation behavior of the two resins. In various commercial polymer processing operations the polymeric resin is subjected to nonhomogeneous stress and temperature fields, re- sulting in distributions of rapidly changing stress and normal stress differences, along with distributions in temperature and pressure, in the melt being shaped. The shaping step also in- volves the rapid cooling of the shaped melt. In such processes, the relaxation and return of the macromolecules to their equi- librium conformation, upon leaving the stress field, is hindered by the quenching. Thus, the relaxation behavior of a resin does influence various properties of articles processed from this resin, including the orientation distribution in the article.

As the predictions of the Wagner model show in Figures 18- 21, there are significant differences in the way the two resins are expected to relax. Figures 18-19 show the relaxation of the shear stress and first normal stress difference following steady shear at 1 to 100 s-l. Polycether imide) relaxes considerably faster than poly(phenylene ether) at the same temperature. It is interesting to compare their behavior at the typical processing

808 KALYON, YU, AND YU KALYON, YU, AND YU

“o,- PpE- pKI -___ 1:j=1

II:+100

Fig. 18. Shear stress relaxation behavior of polycether imide), PEI, and poly(phenylene ether), PPE, following steady shear at 290°C.

=, s

Fig. 19. Primary normal stress relaxation coefficient of polycether imide), PEI, and poly(phenylene ether), PPE, following steady shear at 1 and 100 se’ at 290°C.


-2 I ““““I “f”“‘1 ““““‘1 ““““I ““‘IV 10" m3 10" 10-l lo" ld

TIME,S Fig. 20. Shear stress relaxation behavior of poly(ether imide) and poly-

(phenylene ether) following steady shear at typical processing temperatures.

-5-2 10" lo-" lo-z 10-l lo" lo'

TJMJSS

Fig. 21. Primary normal stress difference relaxation for polycether imide) and poly(phenylene ether) at typical processing temperatures.


temperatures: 330°C for poly(phenylene ether) and 370°C for poly(ether imide). The predicted relaxation of the shearing stress and the first normal stress difference following steady shearing at 100 s-l at these temperatures are shown in Fig- ures 20 and 21, respectively. These results again suggest that the poly(ether imide) relaxes an order of magnitude faster than poly(phenylene ether) under these typical processing conditions as indicated by the time it takes for the torque and the normal force values to decrease to 1% of their initial values. This significant difference observed in the relaxation behavior of these two resins, coupled with the higher first normal stress difference values of poly(phenylene ether) indeed generate appreciable differences in the experimentally determined orientation distributions of specimens injection molded from these two resins.’

CONCLUSIONS

The rheological behavior of poly(phenylene ether) and poly(ether imide) was studied. The relaxation spectra and the damping function data were fitted in terms of a Wagner-type BKZ model for both resins. The predictions of this model were compared with various viscometric material functions determined experimentally in this study. The reported experimental data and the model parameters should be helpful in defining the correct processing windows for these two engineering plastics. The results also have revealed significant differences in the stress relaxation behavior of the two resins, which should have strong ramifications in the microstructure development and ultimate properties of articles processed from these two engineering thermoplastics.

We gratefully acknowledge the financial support of General Electric Company in terms of an unrestricted research grant. We also thank Dr. P. Shenian, Dr. R. Allen, and Dr. L. Schmidt of General Electric for valuable discussions and sug- gestions. Dr. Subir Dey and Mr. Alan Wagner of Stevens prepared some of the vacuum compression-molded samples and carried out some of the characterization work.

References

1. L. Schmidt, J. Appl. Polym. Sci., 23, 2463 (1979). 2. W. Priest and R. Porter, J. Polym. Sci., A-2,10, 1639 (1972).


3. L. Schmidt and J. Emmanuel, Rheology, Volume 2: Fluids, G. Astarita, G. Marrucci, and L. Nicolais, Eds., Plenum, New York, 1980.

4. A. Ausin, I. Equizabal, M. Munoz, J. Pena, and A. Santa Maria, Polym. Eng. Sk., 27, 529 (1987).

5. M. Hansen and D. Bland, Polym. Eng. Sci., 25, 896 (1985). 6. L. Utracki, A. Catani, G. Bata, V. Tan, and M.R. Kamal, J. Appl. Polym.

Sci., 27, 1913 (1982). 7. D. Kalyon, S. Dey, and A. Wagner, “Injection molding of Two Engineering

Plastics: Poly(phenylene ether) and Poly(ether imidel”, Polymer Processing Soci- ety Meeting, Buffalo, NY, September 30, 1987.

8. D. Kalyon, A. Wagner, and S. Dey, “Injection Molding and Microstructure Development of Two Engineering Plastics: Poly(ether imide) and Poly(2,6- dimethyl-1,4-phenylene ether),” submitted to Adv. Polym. Techn., (1988).

9. A. S. Hay, U.S. Patent 3,306,874, General Electric (19671. 10. V. Abolins, D. Aycock, and D. White, Encyclopedia of Polymer Technology,

1987. 11. M. Kryszewski and J. Jachowicz, Developments in Polymer Degradation+!,

N. Grassie, Ed., Applied Science Publishers, London, 1965. 12. S. White, R. Weissman, and R. Kambour, J. Appl. Polym. Sci., 27, 2675

(1982). 13. A. Nadai, Plasticity: A Mechanics of the Plastic State of Matter, McGraw

Hill, New York, 1931 pp. 126-128. 14. R. Penn and E. Kearsley, Trans. Sot. Rheol., 20, 2, 227 (1976). 15. P. Soskey and H. Winter, J. Rheol., 26, 5, 625 (1984). 16. H. Laun, Rheol. Acta, 1’7, 1 (1978). 17. B. Rabinowitsch, Z. Phys. Chem., A145, 1 (1929). 18. J. Dealy, Rheometers for Molten Plastics, Van Nostrand Reinhold, New

York, 1982. 19. B. Bernstein, E. Kearsley, and L. Zapas, Trans. Sot. Rheol., 7, 391 (1963). 20. M. H. Wagner, Rheol. Acta, 15, 136 (19761. 21. A. Lodge, Trans. Faraday Sot., 52, 120 (19561. 22. K. Osaki, “Proceedings of the Seventh International Congress on

Rheology”, C. Klason and J. Kubat, Eds, Tages Anzeiger, Gothenburg, 1976 pp. 1044109.

23. M. Wagner and S. Stephenson, J. Rheol., 23,489 (1979). 24. J. Ferry, Viscoelastic Properties of Polymers, John Wiley, Third Edition,

New York, 1980. 25. W. Cox and E. Merz, J. Polym. Sci., 28, 619 (1958).

Received October 15, 1987 Accepted March 31, 1988

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Melt Rheology of Two Engineering Thermoplastics - Highly Filled