Polymer Blends and Composites in Multiphase Systems - American

11 Rheological Behavior of Blends of Nylon with a Chemically Modified Polyolefin HSIAO-KEN CHUANG and CHANG DAE HAN

Downloaded by UNIV LAVAL on July 11, 2016 | http://pubs.acs.org Publication Date: May 1, 1984 | doi: 10.1021/ba-1984-0206.ch011

Department of Chemical Engineering, Polytechnic Institute of New York, Brooklyn, NY 11201 An experimental study was conducted to investigate the rheological behavior of a heterogeneous polymer blend system consisting of Nylon 6 and a chemically modified polyolefin. The polyolefin is believed to be an ethylene-based multifunctional polymer, although its exact molecular structure has not been disclosed. For the study, phase contrast microscopy was employed for identifying the individual components in the blend; differential scanning calorimetry (DSC) was used for investigating the melting behavior; and a dynamic mechanical analyzer (DMA) was used for investigating the transition behavior of the blends. In the polyolefin-rich blends, the Nylon 6 is the dispersed phase, whereas in the Nylon-rich blends, the polyolefin forms the dispersed phase. The rheological measurements taken do not follow the additivity rule, and show evidence of chemical interaction between the Nylon 6 and polyolefin phases. JL w o TYPES OF P O L Y M E R B L E N D S can be prepared: heterogeneous (i.e., i m

miscible or incompatible) and homogeneous (i.e., miscible or compatible). At a given temperature, homogeneous blends give rise to a single phase i n which individual components are mutually soluble i n one another. T h e properties of these blends usually obey the rule of mixtures; although some times, physical-mechanical properties superior to those of the individual components have been observed. Various experimental techniques, such as electron microscopy, small-angle X-ray diffraction, light scattering, ther mal analysis, and dynamic mechanical analysis, have been used to investi gate the question of miscibility (1,2). The majority of heterogeneous polymer blends contain a matrix phase of rigid resin and a dispersed phase of flexible resin. Naturally, the charac0065-2393/84/0206-0171$06.00/0 © 1984 American Chemical Society

Han; Polymer Blends and Composites in Multiphase Systems Advances in Chemistry; American Chemical Society: Washington, DC, 1984.

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teristics of the blend w i l l be directly influenced by the chemical structure, molecular weight ( M W ) and molecular weight distribution ( M W D ) of each component, and by the ratio of each component i n the blend. H o w ever, knowledge of these variables, although necessary, is not sufficient for the prediction of the ultimate mechanical-physical behavior of polymer blends. They are strongly influenced by many other factors, such as the size, shape, distribution and relative deformability of the flexible dispersed phase, and the nature and extent of adhesion between the phases (3). These factors affect the blend morphology, which is also strongly influenced by the processing conditions (i.e., the deformation and thermal histories). In terfacial adhesion i n heterogeneous polymer blends has a profound influ ence on the mechanical properties of such systems. Basically, the processing of heterogeneous polymer blends involves the following considerations: (1) the rheological properties of the individual constituents; (2) the effectiveness of mixing; (3) control of the microstruc ture i n the solid state; and (4) control of the mechanical-physical properties of the blends. Currently, no rigorous theory can predict w h i c h of the two compo nents i n a blend w i l l form the discrete phase dispersed i n the other compo nent. The state of dispersion (or the mode of dispersion) depends on the rheological properties of the individual polymers, which i n turn are influ enced by their molecular weights (3-6). It would be worth investigating how the phase interactions, if any, might be affected as one polymer or both, present either as the continuous phase or discrete phase, crystallize as the molten polymer blend solidifies during processing. In heterogeneous polymer blends, many interrelated variables affect the mechanical-physical properties of the finished product. For instance, the method of blend preparation (i.e., the method of mixing the polymers and the intensity of mixing) controls the morphology of the blend (i.e, the state of dispersion, domain size, and its distribution), w h i c h , i n turn, con trols the rheological properties of the blend. Currently, no comprehensive theory can predict the mechanical-physical properties of a heterogeneous polymer blend i n terms of its processing variables. Therefore, an investiga tion is needed into the processing-morphology-property relationships of heterogeneous polymer blends. As part of our continuing effort for establishing these relationships, we have recently conducted an experimental investigation w i t h blends of N y lon 6 and an ethylene-based multifunctional polymer (Du Pont C X A 3095). The choice of these blends was based on our earlier experimental evidence (7) that, when the two resins were coextruded forming two stratified lay ers, strong adhesion occurred between the layers. This observation has prompted us to speculate that blends of these two resins (Nylon 6 and C X A 3095) might give rise to some interesting properties. In this chapter, we shall present the highlights of our experimental results.


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Expérimental Materials. The Nylon 6 was a fiber-spinning grade, supplied by American Enka Company. The CXA 3095 (Du Pont Co.) is believed to be an ethylene-based multifunctional polymer; however, neither the chemical composition nor the molecular structure was disclosed to us by the resin manufacturer. We prepared four blends, with the following blend ratios (by weight) : Nylon/ CXA = 80/20; Nylon/CXA = 60/40; Nylon/CXA = 40/60; and Nylon/CXA = 20/ 80. A twin-screw compounding machine (Werner & Pfleiderer) was used.


Sample Preparation. The compounded pellets were compression molded in the form of disks for the rheological measurements. The compression-molded films were also used for dynamic mechanical analysis (DMA), as well as for taking photomicrographs to investigate the state of dispersion in the blend. Thermal and Thermomechanical Analyses. To determine the melting behavior of the materials, we used a Du Pont 910 differential scanning calorimeter (DSC); for measuring the thermomechanical behavior of the materials, we used a Du Pont 981 dynamic mechanical analyzer (DMA) in conjunction with a 1090 thermal analyzer. Before measurements were taken, the samples were dried in a vacuum oven at 65 °C for 48 h, and measurements were carried out in a nitrogen atmosphere at a heating rate of 10 °C/min. Rheological Measurement. A cone-and-plate rheometer (a Weissenberg Model B-16 rheogoniometer) was used to measure steady shearing flow properties at various temperatures. Because Nylon 6 has a melting point of approximately 220 °C and CXA 3095 begins to decompose at approximately 250 °C, the permissible temperature range for rheological measurements (and hence for the melt processing operation) was 220-250 °C. Melt Drawability (Stretchability) Test. To test the drawability of the blends, we conducted a simple extrusion experiment by using a melt indexer equipped with a cooling chamber and a takeup device. The following experimental procedure was employed: The capillary and reservoir sections of the melt indexer were charged with the pellets and heated up to a desired temperature. A known weight was placed on top of the piston rod, and the flow rate was measured. This measurement allowed us to calculate the linear velocity of the melt leaving the die exit. The extrudate, upon exiting from the die, was passed through a cooling chamber into which nitrogen at room temperature was blown gently, and was pulled by a takeup device controlled by a speed controller. The maximum takeup speed was determined by pulling the thread until the thread broke. From this data we determined the maximum draw-down ratio. In the experiment, we used two different lengths of capillary, which gave LID ratios of 4 and 6.3 (D = 2.10 mm). Results and Discussion Morphology of the N y l o n / C X A Blends. T h e compression-molded specimens were first put into microtom capsules filled w i t h a l i q u i d , u n cured epoxy resin. After the epoxy resin was solidified w i t h the a i d of a curing agent, the embedded specimen was microtomed into thin films of


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about 5 μ m thick. The films were then examined under a phase-contrast microscope. Figure 1 shows photomicrographs of four blend samples i n which the dark area represents the Nylon phase and the white area repre sents the C X A phase. Figure 1 shows that, i n the C X A - r i c h blends, the Nylon forms the dis crete phase, dispersed i n the continuous C X A phase, and i n the Nylon-rich blends, the C X A forms the discrete phase dispersed i n the continuous Nylon phase. The blend ratio appears to have played a predominant role i n deter mining which of the two components forms the discrete and w h i c h forms the continuous phase. Melting Behavior of the Nylon/CXA Blends. Figure 2 displays ther mograms of the blends investigated. Nylon 6 has a melting point of 222 °C, and C X A has two melting points, 126 and 98 ° C . The Nylon/CXA = 20/80 and Nylon/CXA = 40/60 blends have three melting points (222, 126, and 98 ° C ) , whereas the Nylon/CXA = 80/20 and Nylon/CXA = 60/40 blends have two melting points (222 and 98 ° C ) . N o depression of the melting point occurred i n either the Nylon-rich or C X A - r i c h blends, except that i n the Nylon-rich blends (i.e., Nylon/CXA = 80/20 and Nylon/CXA = 60/40)

Figure 1. Photomicrographs of the compression-molded specimens of the Nylon/CXA blend system: (a) Nylon/CXA = 20/80, (b) Nylon/CXA = 40/ 60, (c) Nylon/CXA = 60/40, and (d) Nylon/CXA = 80/20.



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0

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Temperature (°C)

Figure 2. DSC thermograms of the Nylon/CXA blend system: (1) Nylon, (2) Nylon/CXA = 80/20, (3) Nylon/CXA = 60/40, (4) Nylon/ CXA = 40/60, (5) Nylon/CXA = 20/80, and (6) CXA 3095.

the secondary melt point (126 °C) of the C X A disappears completely i n the thermogram. Because, at present, we do not know the exact molecular structure of the C X A , we cannot relate the individual melting point with a particular structure. Nevertheless we know that the C X A is a blend of two polymers, one of which is a copolymer. Therefore, two distinct structures i n the C X A must be responsible for the existence of two melting points. Of great inter est is the disappearance of the second melting point (126 °C) i n the thermograms of the Nylon-rich blends. O n the basis of the photomicro graphs shown i n Figure 1, the disappearance of the secondary melting point (126°C) i n the thermogram occurs only when the C X A forms the dis crete phase. The reason for this seemingly peculiar melting behavior is worth investigating i n the future. Dynamic Mechanical Behavior of the Nylon/CXA Blends. Figure 3 presents plots of the loss modulus (E ") of the blends investigated. Pure N y lon has three peaks: a peak at 64 °C, β peak at - 50 °C, and y peak at - 1 2 0 °C, although the y peak is not shown i n Figure 3. The a peak is believed to be associated w i t h the glass transition, the β peak is attributed to the existence of the polar group forming the hydrogen bonds i n the polyamide, and the y peak is thought to be associated w i t h the crankshaft rotation of the - ( C H ) group i n the main chain of the polyamide (8). 2

n



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Figure 3 shows one peak at - 25 °C for C X A 3095. Because the molec ular structure of the C X A 3095 is not known, we cannot comment on which type of functional group i n the C X A 3095 is responsible for the peak. Also, in the Nylon-rich blends, the a transition of Nylon is shifted slightly toward lower temperatures i n the Nylon/CXA = 80/20 and Nylon/CXA = 60/40 blends, and the β transition of Nylon is shifted somewhat toward higher temperatures ( - 38 and - 32 °C i n the Nylon/CXA = 80/20 and Nylon/ C X A = 60/40 blends, respectively). The rather noticeable shift of the β transition of Nylon i n the Nylon-rich blends seems to indicate some kind of chemical interaction between the C X A 3095 and the N y l o n 6. In the C X A - r i c h blends, as shown i n Figure 3, the peak ( - 25 °C) of the C X A 3095 is shifted slightly toward lower temperatures i n the Nylon/ C X A = 20/80 and N y l o n / C X A = 40/60 blends, and no peak representing the Nylon 6 is observed. Therefore, the N y l o n 6 has little influence on the transition behavior of the C X A - r i c h blends. Figure 4 displays plots of the elastic (Young's) modulus (E ' ) of the blends investigated. The Ε ' decreases as the Nylon content i n the blend is decreased, and at temperatures of 50 °C and higher, the Nylon 6 i n the C X A - r i c h blends contributes little to the modulus of the C X A 3095. Rheological Behavior of the Nylon/CXA Blends. Figure 5 gives plots of viscosity η versus shear rate y, and Figure 6 shows plots of normal stress

0.28 h

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Figure 3. Loss modulus vs. temperature for the Nylon/CXA blend system: (1) Nylon 6, (2) Nylon/CXA = 80/20, (3) Nylon/CXA = 60/40, (4) Nylon/ CXA = 40/60, (5) Nylon/CXA = 20/80, and (6) CXA 3095.


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11.

Figure 5. Viscosity vs. shear rate for the Nylon/CXA blend system at vari ous temperatures (°C): (1) Nylon—Θ, 230; Φ, 240; 3 , 250; (2) Nylon/CXA = 80/20— •, 230; •, 240; Œ, 250; (3) Nylon/CXA = 60/40—ψ, 230; Τ , 240; Ψ, 250; (4) Nylon/CXA = 40/60—φ, 220; 0 , 230; * , 240; (5) M/Z 60

ι 70

ι 80

ι 90

I 100

(wt%)

Figure 10. Viscosity vs. blend ratio for the polypropylene (PP)/polystyrene (PS) blend system (T = 200 °C) at two different shear stresses (N/m ) (11). Key: © , 4.1 Χ 10 ; and A , 4.8 Χ 10 . 2

4

4


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Melt Drawability of the Nylon/CXA Blends. Figure 11 displays the observed melt drawability (or stretchability) of the N y l o n / C X A blends i n a simple extrusion experiment. The maximum draw-down ratio (V IV ) used in Figure 11 may be considered to be a simple, effective test for deter mining how much more stretchable one material is than others, if the tests are conducted under the same extrusion conditions. Figure 11 demon strates two important points. First, the (V I V ) of the C X A - r i c h blends depends on deformation history (i.e., L/D ratio), whereas the (V /V ) of the Nylon-rich blends does not. This result is understandable because the elasticity of the C X A - r i c h blends is greater than that of the Nylon-rich blends, as may be seen i n Figure 6. Second, except for the N y l o n / C X A = 40/60 blend, the blends give rise to a melt drawability better than that ex pected from the additivity rule. W h y the Nylon/CXA = 40/60 blend gives rise to such a poor melt drawability is inexplicable from the experimental evidence. L

L

0 MAX

L


0 MAX

0 MAX

Concluding Remarks From the observations on the Ε ' of the blends, given in Figure 4, the N y lon-rich blends (e.g., Nylon/CXA — 80/20 blend) seem to be more attrac•0*F=

0

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10

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Nylon ( wt

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Figure 11. Maximum drawn-down ratio vs. blend ratio for the Nylon/CXA blend system (T = 255 °C), extruded through two different capillaries. Key: O, L/D = 4.0; and Δ, L/D = 6.3.


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tive than the C X A - r i c h blends. The C X A has a higher viscosity than the Nylon (see Figure 6) and therefore, when subjected to a flow field, the C X A w i l l not deform as easily as the Nylon. The Nylon-rich blends retain essen tially the same melting and transition behaviors as those of Nylon (see Fig ures 2 and 3). W e speculate that, during melt processing, some chemical interactions occurred at the interface between the C X A and Nylon phases and gave rise to unique properties (e.g., melt draw ability) i n the Nylonrich blends. In the future, we shall put our efforts into unraveling the mechanism(s) of chemical interaction that may exist i n Nylon/CXA blends.


Acknowledgment W e acknowledge, with gratitude, that American E n k a Company supplied us with the large quantity of the Nylon 6 resin, and that the D u Pont C o m pany supplied us with the C X A 3095. Literature

Cited

1. "Polymer Blends"; Paul, D. R.; Newman, S. Eds.; Academic: New York, 1978. 2. Olabisi, O; Robeson, L. M.; Shaw, M. T. In "Polymer-Polymer Miscibility"; Academic: New York, 1979. 3. Han, C. D. "Multiphase Flow in Polymer Processing"; Academic: New York, 1981. 4. VanOene, H. J. Colloid Interface Sci. 1972,

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5. 6.

Han, Han,

C. D . Yu, T. C. J. Appl. Polym. Sci. 1971, C. D.; Kim, Y. W. Trans. Soc. Rheol. 1975,

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C. D.; Lem,

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15, 19,

1163. 245.

7. Kim, Y. J.; Han, C. D., Unpublished research, 1981. 8. Kawaguchi, T. J. Appl. Polym. Sci. 1959, 2, 56. 9. Han, C. D. In "Rheology in Polymer Processing"; Academic: New York, 1976. K. W. Polym.

Eng.

Reviews 1982, 2, 135. Appl Polym. Sci. 1975, A 1932, 138, 41.

11. Han, C. D.; Kim, Y. W.; Chen, S. J. J. 12. Taylor, G. I. Proc. R. Soc. London, Ser. RECEIVED

for review January 20, 1983.

ACCEPTED

August 29,

19,

1983.


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Polymer Blends and Composites in Multiphase Systems - American

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