Service Life Prediction - American Chemical Society

Chapter 13

Downloaded by NORTH CAROLINA STATE UNIV on September 17, 2012 | http://pubs.acs.org Publication Date: November 21, 2001 | doi: 10.1021/bk-2002-0805.ch013

Mechanical Properties of Painted TPO Plastic Resulting from Morphological Interphase Variations Rose A. Ryntz Visteon Corporation, 401 Southfield Road, Dearborn, MI 48121

The physical behavior of painted and/or unpainted thermoplastics subjected to applied stress in field applications varies dramatically depending upon material selection and processing parameters. Of particular concern when trying to relate chemical structure of the paint/plastic to end-use properties is the "interphase" miscibility within the plastic alloy. In this paper, the "interphase" management of elastomer, as dispersed within a poly(propylene) matrix, is discussed and related to resultant surface damageability, e.g., scratch (of unpainted specimens) and friction induced paint damage (a compressive shear loading event) of painted specimens. The role of polymer processing, in particular injection molding shear velocity, as well as the paint process conditions utilized, on the "interphase" between the elastomer/poly(propylene) matrix, will be discussed. It was determined that by controlling molecular weight, molecular weight distribution, crystallinity, and melt viscosities of the elastomer/poly(propylene) matrix, surface damageability of the fabricated part could be controlled.

In injection molded plastics, in particular where semi-crystalline polymers are involved, residual stresses can occur in the top few microns of the

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© 2002 American Chemical Society

In Service Life Prediction; Martin, J., et al.; ACS Symposium Series; American Chemical Society: Washington, DC, 2001.


267 plastic depending upon injection molding conditions and polymer composition utilized [/]. It was reported by Schonhorn and others [1-7] that heterogeneous nucleation and crystallization of polymer melts against high-energy surfaces, e.g., metals, metal oxides, and alkali halide crystals, results in marked changes in the surface region morphology. More specifically, it was claimed that transcrystallinity [8,9] (transcrystallinity is reported to consist of elongated spherulites originating at the polymer surface and propagating for several microns normal to the surface) could be induced in polyethylene and polypropylene by crystallizing the respective melts in intimate contact with specially prepared aluminum and copper. Formation of a transcrystalline layer is favored by conditions that induce a high density of nuclei at the surface. A close arrangement of growth centers causes the spherulites to grow in a columnar fashion with little lateral development [10]. Microstructures of semi-crystalline polymers vary from spherulitic to lamellar and single crystals, in size of structural units, depending upon [77]: • • • •

molding temperature low patterns aging and heat treatment nucleating agents

The distribution of spherulite sizes varies as a function of depth. In injection molded polypropylene, for example, not only does the crystal size vary but also the crystal type. Hexagonal spherulites (less perfect flat, concave boundaries) were shown to be more prominent at the surface. They were characterized by greater susceptibility to etching in SEM analysis. Monoclinic spherulites, a more crystalline, ordered arrangement, were found to occur in the bulk morphology in conical shapes. In thermoplastic olefin (TPO, a blend of polypropylene and ethylene propylene diene rubber (EPDM)) that had been injection molded, Bonnerup [72] found that not only was the polypropylene surface crystallinity affected but that the EPDM phase separated in the crystallization process and was present in very low concentrations at the surface. Bikkerman [73] found that two solids in contact couldn't fail exactly at the interface between them. Hence, if failure occurs at or near an interface at a relatively small-applied stress a weak boundary layer is assumed to have been present. Weak boundary layers are believed to arise if: • • •

there is an incompatibility between two polymers so that they remain separate; the surface roughness amplitude between two similar solids is 3000 angstroms or less; and a similar Tg for both polymers exists so that annealing eliminates shrinkage and thermal stresses. Fracture toughness


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and crack propagation are related to weak boundary layer management. Injection molding is one of the most widely utilized methods of polymer processing. This process includes injecting molten polymer into a cold mold followed by packing under high pressure and subsequent cooling to solidification. The filling and cooling stages have an important effect on the rheological properties since the viscoelastic nature of the polymer results in development of shear and normal stresses and large elastic deformation during filling with subsequent incomplete relaxation during the cooling stage. The resultant residual stresses, which determine the orientation in the final molded part, are dependent upon the thermal, rheological, and relaxation properties of the polymeric material as well as the processing conditions [14]. Molecular orientation in injection molding has been modeled by many researchers [75, 16] to explain the complex molecular orientation distribution observed. Most models incorporate flow and heat transfer mechanisms coupled with molecular theories. The orientation in the surface skin is related to steady elongational flow in the advancing front, whereas the orientation in the core is related to the shear flow, behind the front, between two solidifying layers [75]. Coupled with the elongational and shear-induced orientations, a molecular relaxation process takes place that is determined by the rate of heat transfer. Internal stresses that develop within the injection-molded part are the result of thermal, flow, and pressure histories[7 7]. The melt temperature of the polymer was found to cause two maxima in residual stress (R.S.) [78]. The second one reverses from compressive to tensile. In general, most changes occur in the surface regions, while R.S. decreases with increasing melt temperature, as is the case in zones far away from the gate [19]. It was found that R.S. are compressive in the surface layers and tend to decrease upon increase in mold temperature and distance from the gate. Residual stress can be measured by optical birefringence. Birefringence has a number of causes. The polarizability of chemical bonds change when they are stressed, giving rise to the photo-elastic response upon which measurements of residual stress may be measured. On the other hand, chemical bonds are directional and in a non-isotropic material the presence of favored bond orientations will produce birefringence, so that in an injection molded article containing frozen in molecular orientation this may provide a much larger contribution than the residual stresses [20]. The birefringence effect should be higher closest to the mold wall, dependent of course upon the heat diffusivity of the polymer and the molding conditions. Injection molding conditions, e.g., melt temperature, aging and heat treatment, surface geometry, etc., have been shown to be significant in



269 determining the final cohesive strength of a TPO substrate [21]. The crystallinity and phase separation that occurs in a TPO sample once molded can be controlled to some degree based on the molding history. The degree of cohesive strength in a TPO sample, however, seems to be a direct result of the molecular weight, molecular weight distribution, and miscibility of the blends of elastomer and PP that are chosen. In this paper, an attempt is made to relate structure of the TPO blend chosen to surface damageability caused by scratching, and friction induced compressive shearing ("gouging"). The role of interphase management, e.g., control of miscibility between alloying agents, appears to be the major factor affecting the ability of the plastic part to resist surface damage caused by external forces.

Data and Results The TPOs evaluated in this study were made by physically dry-blending the components (Table 1) at 20 wt% elastomer in 80 wt% PP (where filler was utilized it was done so at 10 wt%) and melt extruding the blend through a Werner-Pfleiderer twin screw extruder using a general compounding profile and a single strand die. The extruder was preheated to the following barrel conditions: 210°C, 220°C, 220°C, 220°C, 220°C, 225°C. The screw was run at a constant speed of 115 rpm. Poly(propylene) and elastomer were obtained from Exxon Chemical, Bayport, TX or Dow-DuPont Elastomers, Midland. MI. Melt flow rates (MFR) of each component were measured according to ASTM D1238-96, 230/2.16. Melting points and percent crystallinities were determined on a DuPont Model TA Modulated Differential Scanning Calorimeter. Molecular weights and molecular weight distributions were supplied by Exxon Chemical and DuPont Dow Elastomers. Talc (Cimpact) was obtained from Luzenac Minerals, Denver, CO. Wolastonites were obtained from Nyco Minerals, Inc., Charlottesville, VA. The supplier performed all particle size measurements. Each blend was molded into plaques on a Cincinnati Milacron 110 ton. 5 ounce injection molding machine at injection velocities of 1.27 and 7.62 cm/sec. Tensile testing on unpainted plaques was performed on an Instron Model 6025 tensile tester equipped with a 454.5 kg load cell and Instron series IX computer controlled software. Work of fracture data was calculated from the total area under the stress-strain curve. Sample specimens utilized in the work of fracture calculations were 5 χ 17.8 cm bars modified with varying ligament lengths inscribed through use of double-sided notches obtained with a


270 razor blade. Essential work values were calculated according to the method described by Wu [22]. Flexural modulus (ASTM D790-96) and Izod impact (ASTM D256-93, method A) data were measured on unpainted samples. Plastic moduli (H ias) and degree of plastic deformation (plasticity) (W /W ) were measured on unpainted specimens with a Fisherscope H100 with a load of 1000 mN utilizing a Vickers indentor. Optical microscopy analysis of 30 micron thick unpainted specimens, obtained by cryogenically microtoming the specimens (direction of cut was parallel to the flow direction in the plaque) using Histoprep media, was measured on a Leica microscope equipped with cross-polarizing filters. The samples were mounted in Canada balsam between two microscope slides prior to analysis. Depth of boundary layers within the specimens was measured utilizing Optimus optical imaging software. Scanning electron microscopy (SEM) was performed on selected samples on cryogenically fractured surfaces of TPO under a magnification of 1500x. Scratch testing was performed according to the Ford Laboratory Test Method (FLTM) BN 108-13. In this FLTM, the panel is placed on a movable platen onto which is placed a beam containing a scratch pin. The beam is 250 mm long and is equipped with a scratch pin that consists of a highly polished steel ball (1 mm +/- 0.1 mm in diameter). The beam is loaded with a weight ranged of 7.0 N. The beam is driven by compressed air to draw the pin across the surface of the plaque to generate a scratch. Sliding velocity was maintained at approximately 100 mm/sec and all tests were performed at 25°C. The samples were conditioned at 25°C for 24 hours prior to scratching. Scratched samples were analyzed with an interferometer at intervals of 1 hour and 24 hours after scratching. Scratches produced ranged from 1 to 3 μιη in depth depending upon the plaque evaluated. Scratch deformations are reported as depth of deformation (in microns) as referenced to the unscratched surface. Panels were painted through use of an adhesion promoter (7.5 μπι dry film thickness of a chlorinated polyolefin primer), white basecoat (37.5 μπι dry film thickness, a one-component (IK) melamine crosslinked basecoat or a twocomponent (2K) urethane basecoat, and a 2K urethane clearcoat (50 μπι dry film thickness), which were applied wet-on-wet and subsequently baked for 30 minutes at 121°C ambient. Friction induced paint damage resistance ("gouge") was measured on either a Ford proprietary friction induced paint damage apparatus (STATRAM) or a commercially available FIPD device (SLIDO) utilizing a 2 design of experiments (DOE), varying parameters such as sliding velocity, temperature, acceleration, and compressive (vertical) or traction (horizontal) force [23, 24]. Results are presented as either the area of gouge damage or as the percentage of r


p

3


t

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panels that failed in the painted TPO FIPD testing obtained from the design of experiments (3 replicates under each experimentally designed run were measured). Table 1 lists the physical properties of the resins utilized to make the TPO blends in this study. As can be seen in the Table, the elastomer types were chosen to reflect varying solubility characteristics with the PP as well as varying crystallinity and melt flow rates. The following studies will look at the influence of melt flow ratio of elastomer/PP in the TPO samples as well as the effect of elastomer crystallinity on physical properties attained.

Table 1 Properties of TPO Components Polymer*

Density (g/crrf)

1042 PP JSR07P 4033 EB 3125EB 3022 EB 8150 EO 8180 EO

0.9049 0.8589 0.8837 0.9124 0.9057 0.868 0.863

MFR DSC Melt % (dg/min) CO Crystallinity

1.9

Service Life Prediction - American Chemical Society

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