The Morphology and Mechanical Properties of Ethylene-Propylene

19 The Morphology and Mechanical Properties of Ethylene-Propylene Copolymers and Blends Downloaded by COLUMBIA UNIV on March 15, 2013 | http://pubs.acs.org Publication Date: December 9, 1985 | doi: 10.1021/ba-1986-0211.ch019

P. P R E N T I C E , E . P A P A P O S T O L O U , and J. G . W I L L I A M S Department of Mechanical Engineering, Imperial College of Science and Technology, L o n d o n , England S W 7 2 B X

The mechanical properties of the group of materials usually referred to as ethylene-propylene block copolymers were found to resemble those of physical blends. According to optical and scanning electron microscopy, an excluded phase exists in these materials that is rub ber-like in nature and has dimensions in the range of 0.7 to 5 μm, depending upon the source of supply. Simple extraction experiments demonstrated that the majority of this phase is removable by physi cal means (i.e., it is not chemically bound to the polypropylene ma trix). The loss factor, tan δ, plotted as a function of temperature, resulted in two peaks in both the copolymers and the blends. One peak corresponded to the matrix, and the other peak corresponded to the ethylene-rich phase. From the limited data so far available, we found a correlation between the size of the tan δ peak and the ethylene content.

THE RELATV IELY LOW COST

and general availability of polyolefins such as polyethylene (PE) and polypropylene (PP), along with their many and var ied properties, have created a broad market for these materials. However, not all the properties of these materials are beneficial, for instance the sus ceptibility of high-density polyethylene ( H D P E ) to environmental-stress cracking and the relatively high glass transition temperature (T g ) of PP, which makes it unsuitable for low-temperature applications. Therefore, polyolefin development has concentrated on improving the performance of existing materials. T w o areas that have been investigated are the produc tion of physical blends of the homopolymers with various types and amounts of thermoplastic rubbers and the development of copolymers. T w o types of copolymers are available commercially. T h e first type is random copolymers in which, as their name implies, the comonomer is in corporated into the polymer chain at irregular intervals—this group in0065-2393/86/0211/0325$06.00/0 © 1986 American Chemical Society

In Multicomponent Polymer Materials; Paul, D., et al.; Advances in Chemistry; American Chemical Society: Washington, DC, 1985.

Downloaded by COLUMBIA UNIV on March 15, 2013 | http://pubs.acs.org Publication Date: December 9, 1985 | doi: 10.1021/ba-1986-0211.ch019

326

MULT C IOMPONENT POLYMER MATERA ILS

eludes the ethylene-propylene thermoplastic rubbers (EPR). T h e second group contains the so-called "block" copolymers in which the ethylene comonomer is assumed to be concentrated in distinct regions of the chain. In general, the term block copolymer applies to materials in which the various components of the polymer chain are chemically linked. In this chapter, we hope to show that the ethylene-rich regions in commercially available "block" copolymers of ethylene and propylene are not chemically linked to the matrix. Thus, although these materials are copolymers, it may not be technically correct to refer to them as block copolymers. T h e problem arises because of the considerable uncertainty that exists regarding the true architecture of these materials; most have been assigned structures deduced almost entirely from the synthetic procedure employed. However, the simplest manufacturing process involves the consecutive introduction of two monomers into a reaction vessel, and the existence of residues of the first monomer in the presence of the second leads to a situation in which the relative monomer concentrations are unknown and constantly changing. T h e dynamics of a Ziegler-Natta catalyzed reaction are such that the mean lifetime of growing polymer chains at 70 ° C is about 10 to 16 m i n (J). After this period of time, they are terminated by a chain-transfer mechanism (2). Because, in some processes, the interval between the start of the reaction and the introduction of the second monomer is much greater than this lifetime, many of the initial chains will terminate before encountering the second monomer. For ethylene-propylene copolymers in which propylene is the initial monomer, the addition of ethylene as the comonomer will lead to the initiation of many new chains with an ethylene molecule, because of its smaller size and greater reactivity, while still in an environment rich in propylene. Thus, a large proportion of the ethylene-rich segments will not be chemically linked to the PP preblock. For these reasons, during the transition from one monomer feedstock to another, the resulting product is likely to contain large proportions of an E P R . T h e PP and the rubber are mutually incompatible in the melt (3), and the rubber forms a dispersed phase when the sample is cooled. Optical (4, 5) and scanning electron microscopy (SEM) provide direct visual evidence of this two-phase morphology. If the presence of a second, more rubbery, phase is assumed, dynamic mechanical studies should result in the detection of more than one transition in copolymer systems.

Experimental Materials. Several samples of copolymers from various manufacturers, a l l commercially available, were examined along w i t h some physical blends. The total ethylene content of these materials ranged from 3 to 14 w t % . Microscopy. T h e preparation of specimens and the techniques of optical m i croscopy are described elsewhere (4-6).


19.

PRENTC IE ET AL.

Ethylene-Propylene

Copolymers and Blends

327


For investigation w i t h an S E M (Joel JSM-T200), optically smooth surfaces were prepared from molded plates by using a freshly prepared glass knife. T h e specimens were then treated w i t h p-xylene i n an ultrasonic shaker for about 5 m i n to remove any soluble material from the surface regions. Surfaces treated i n this way are superior (3) to those prepared by acid etching, and they are indistinguish able from those prepared by solvent etching i n a laboratory shaker for 1 to 24 h . The prepared surfaces were subsequently coated w i t h gold by using an Edwards sputter-coating device before viewing i n the S E M at an excitation voltage of 15 k V . Impact Testing. C h a r p y impact specimens were cut from water-quenched plates molded at 190 ° C . Specimen dimensions were 50 x 5.5 m m w i t h a 6 0 ° notch of root radius 15 μπι introduced by means of a fly cutter. Specimens w i t h different notch depths were tested on a pendulum machine, the principles of w h i c h were previously reported by Casiraghi (7), at temperatures ranging from - 80 to + 20 ° C . By using this approach, a fracture-mechanics analysis may be used to deter mine the strain-energy release rate, or impact fracture toughness, Gc. Such an analysis was originally developed (8) for brittle polymers, but was extended (9) to more ductile materials. Gc is related to the fracture energy, the specimen and notch dimensions, and the compliance by the equation W = GcBD + Wk where W is the fracture energy; β and D are the specimen thickness and depth, respectively; and φ is the compliance calibration factor (a function of the specimen geometry a n d notch dimension) that may be obtained from tables (9). is the kinetic energy possessed by the specimen as it is thrown forward after fracture. A plot of fracture energy versus ΒΌφ yields a straight line w i t h a slope of G and a positive intercept on the energy axis w i t h the value of W * . Specimens for dynamic mechanical testing were thin films that were also com pression molded at 190 ° C and water quenched. The specimens, approximately 30 m m long and 4 m m wide w i t h a nominal thickness between 0.15 and 0.25 m m , were examined on a direct-reading viscoelastometer (Rheovibron D D V - 1 1 ) operat ing at 11 H z b y using a simple tensile stress system. c

Results and Discussion Figure 1 shows a typical copolymer viewed by differential interference contrast (DIC) in an optical microscope and displays the characteristic spherulite structure of polyolefin. W i t h i n the bulk of the spherulite, small inclusions are clearly visible that demonstrate the presence of a second phase. Preliminary experiments (I) in these laboratories have demonstrated that this second phase may be extracted, to a greater or lesser extent de pending on the origin of the sample, by refluxing in η - h e p t a n e in a Soxhlet extraction apparatus. T h e P E content of the extracted sample, determined by IR spectroscopy (II), decreases substantially after extraction. T h e mod ulus of the extracted rubber may also vary from one copolymer to another, and this variation may have a profound influence on the mechanical prop erties of the copolymer. This situation is reflected by the electron micro graphs of the freshly microtomed surfaces of a polymer blend and two co-



328


Figure 1. Optical micrograph

(viewed in DIC) of α 5-μηι section of a copolymer.

polymers after treatment with p-xylene in an ultrasonic shaker for 5 m i n at room temperature. In a ternary blend of P P , E P R , and H D P E (85:10:5) (Figure 2) and in one of the copolymers (Figure 3), the resulting cavities are clean and devoid of any included matter. In another copolymer from a dif ferent supplier, however (see Figure 4), a significant amount of debris re mains even after prolonged exosure to p-xylene. This difference between copolymers is evidence that, in the blend and one of the copolymers, the dispersed phase is completely separate and has no chemical linkages to the PP matrix. In the second copolymer, however, a significant amount of true block copolymer acts, in effect, as an emulsifier and cannot be removed from the matrix by simple physical means. T h e only explanation for the differences between the two copolymer samples appears to relate to the manufacturing methods employed. Very few of the materials examined af ter extraction displayed the phenomenon of cavity residues. Therefore, most ethylene-propylene copolymers of this type available commercially must exist as simple blends of homopolypropylene and thermoplastic rub ber (with the possibility, in some systems, of a third component—linear PE). T h e rubber component of a ternary blend may completely envelop (3) the H D P E and prevent the latter from cocrystallizing with the PP matrix. The major advantages of a copolymer over a physical blend are a marked improvement in the dispersion of the secondary phase in the copolymers, even after extensive compounding of the blends in a Banbury-type mixer,



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Ethylene-Propylene

Copolymers

and Blends

Figure 2. SEM of a microtomed surface of a blend (85 % PP, 10 % EPR, and 5 % HDPE) extracted with p-xylene.

Figure 3. Electron micrograph of an extracted microtomed copolymer (A).

surface of a


329


330


Figure 4. Electron micrograph of an extracted microtomed second copolymer (B).

surface of a

and the smaller size of the dispersed phase in many of the copolymers. In their work on blends, Stehling and coworkers (3) found that the optimum diameter of the dispersed phase for toughening PP was less than 0.6 μιχι. The actual size has a marked effect on the impact strength. T h e impact and dynamic mechanical data of a selection of copolymers and blends are plotted on the same set of axes in Figures 5a-5d; each plot corresponds to one material. A l l curves, either for a copolymer or blend, are very similar, and only small variations are apparent in the parameters described. T w o major transitions are evident in the temperature range in vestigated. They are manifested as two peaks in the tan δ versus tempera ture plot, one at about - 40 to - 50 ° C and another, larger peak at about + 10 ° C ; they also appear as deviations in the in-phase modulus (G ') versus temperature plot occurring at about the same temperatures. T h e main transition at about 10 ° C corresponds to the homopolypropylene. T h e of ten-quoted (12) Τg of PP, a result of main-chain relaxations of the amor phous regions of the solid, is around 5 ° C . T h e corresponding T g of P E is less well defined, but the increased flexibility of the P E chain compared with PP would be expected to lead to a significantly lower value. One value quoted with some confidence (12,13) is - 20 ° C . Transition temperatures are highly rate dependent, and the shift in the value for the T g of P P , from 5 to about 10 ° C , is because of this dependence. Values quoted for T g are often evaluated on a torsion pendulum at a frequency of 1 H z , while the data presented here were obtained at 11 H z on a Rheovibron. Even allow ing for this shift, the peak at about - 45 ° C is unlikely to be due to linear



Figures 5a and b. G ' , G c , and tan δ plotted as a function of temperature:

a, copolymer C; and b, physical


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The Morphology and Mechanical Properties of Ethylene-Propylene

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