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Evolution of the microstructure and mechanical properties of polyamide 6 during impact.


The impact resistance of polymers determines their usefulness in many applications. Consequently, there is an increasing demand for tests that accurately simulate end-use conditions. In the past, the most commonly used tests were Charpy and Drop tests. Despite the fact that these tests have been recently instrumented and now generate force versus displacement data, they do not yet provide an intrinsic characterization of the material. Even from a practical view, it is often difficult to unambiguously determine the exact yielding of the polymer, the exact initial crack or origin of rupture.

Additionally, many engineering thermoplastics exhibit a large local plastic deformation even during impact. For example, this has been reported for the Dart impact technique (impact of a striker on a plaque) on various polymers such as polycarbonate (1), polypropylene copolymers (2-4), linear low-density polyethylene (5) or polyamide 12 (6), using either a hemispherical (1, 3-6) or a fiat-headed striker (2). These highly inhomogeneous deformations are almost impossible to characterize or to predict without the use of accurate numerical simulations.

Numerical calculations are nowadays relatively efficient but need, as a first step, an accurate characterization of the behavior of the polymer, i.e., its constitutive equation. A dart test is often used as a tool for the validation of both the code and the constitutive equation (1, 4, 6, 7-10). In this field, calculations are in quite good agreement with experimental observations from a qualitative view. The location of the deformation, the shape of the force versus time curve are often well modeled. Quantitative agreement can be obtained at low deformation (i.e., up to the initiation of the plastic deformation) using more or less simple constitutive equations, e.g., Symonds-Cooper's law (7), bi- or tri-linear laws (1, 8, 9) or more sophisticated viscoplastic approaches (4, 6, 10) based on the G'sell-Jonas' law (11).

Nevertheless, large plastic flows are not modeled in a satisfactory way. In most cases calculations use, as input data, parameters obtained during laboratory tests, e.g., isothermal tensile tests performed at low strain-rates. To extrapolate measurements to impact situation, effects of temperature, strain-rate, strain-hardening and elastic recovery have to be accounted (10). It has been pointed out, using simple linear laws (1, 8), that strain-hardening processes in polymers are of prime importance to model puncture even at low strain-rate. In fact, strain-hardening governs the initiation and the propagation of necking even during multiaxial loadings. More recently, using an equivalent model, but a more accurate constitutive equation (4, 6), it was found that rheological parameters determined during tensile tests are not appropriate to model puncture or impact, even if they were determined in the same range of strain-rate and temperature. It was confirmed that the main physical parameter to be considered to explain discrepancies is strain-hardening, which is largely overestimated in tensile tests with respect to impact tests (4, 6). As strain-hardening is a trace of the microstructural evolution within the material, it was suggested that it was not surprising that uniaxial and multiaxial loadings led to different results. Consequently, one of the main problems concerning the modeling of impact loading remains the anisotropy induced by the deformation itself. This kind of problem has already been mentioned, even in more simple loading situations [e.g., compression (12)].

These observations are the starting point of this study. As it is established that induced anisotropy is a key parameter for the understanding of impact behavior, it is necessary to first characterize it. Important questions to be answered are: how is the deformation initiated and how does it propagate? To solve this problem attempts have been made in the past using different methods:

* high-speed video or photography techniques (13),

* partial penetration at controlled depths (3),

* partial penetration at controlled energies (5).

High-speed video or photography are often difficult to use, owing to cost of the system and the geometry of the experimental devices. Controlled-depth tests, using a hydraulic machine, cannot be done at high velocities, since it is still difficult to stop a high-speed piston at controlled displacements. Partial penetration at controlled energy allows measurements at higher velocities, closer to actual impact situation, but their main disadvantage is that velocity is not constant during the tests. Despite this drawback we chose to use this technique in this paper. Previous works (3, 5) mainly dealt with the description of the shapes of the impacted samples and force versus time data. This study aims to be more complete since the plastic deformation of a polyamide 6 has been quantified as a function of impact energy, in parallel with its microstructural and mechanical evolution. To achieve that point, partial and total Dart tests, using a hemispherical striker were performed. The force versus displacement curves were correlated to the description of the samples including:

* shape description,

* thickness distribution,

* crystalline orientation and microstructure,

* mechanical anisotropy.

The general purposes are to visualize the exact local deformation of the polymer on one hand and to characterize the resulting mechanical properties of the deformed polymer on the other hand, that is, to improve the knowledge of the Induced anisotropy and its effects on the behavior of the polymer.


Procedure for Impacts

An Ultramid B3 polyamide 6 (supplied by BASF), with average molar mass [M.sub.n] 18,000 g/mol, was used. Its peak melting temperature and viscosity were 220 [degrees] C and 140 Pa.s at 250 [degrees] C and 1000 [s.sup.-1], respectively. Samples were cut in disk-shaped injection moulded parts. The diameter of the disks was 182 mm and their thickness 2 [+ or -] 0.14 mm. Injection was performed using a 300 t injection machine. The temperature of the polyamide was 250 [degrees] C and the temperature of the mold 44 [degrees] C. The injection velocity was 60 mm/s. The packing time was 40 s.

These injection conditions led to parts whose microstructure is quite homogeneous (i.e., from one point of the disk to another). Optical microscopy on thin slices (8 [[micro]meter] thick) reveals in the thickness a skin to core microstructure which can be distinguished by three zones:

* a [approximately]50 [[micro]meter]-thick skin zone in which the crystalline morphology cannot be observed,

* a shear zone of equivalent thickness in which spherulites of increasing sizes are observed,

* a core zone in which spherulites with a diameter from 20 to 30 [[micro]meter] are present.

Wide-angle X-ray diffraction using the fiat film camera technique (filtered Cu K[Alpha] radiation) does not reveal any crystalline orientation in the disks.

Square samples of 50 mm width were cut in the disks. They did not include the injection gate. Impact tests were performed using screwed clamping devices, which defined a circular impact zone with a 40 mm diameter. The striker was a cylinder with a hemispherical head whose diameter was 20 mm.

Partial impacts of increasing energy were performed with a falling mass apparatus. This technique permits the observation of the successive stages of the deformation during an impact. The device used had a fixed striker in its lower part. The force on the striker was measured as a function of time thanks to a piezoelectric sensor. The sample was firmly screwed on a 16.8 kg trolley falling from different heights ranging from 4 to 184 cm. As a consequence, impact energy and initial velocity ranged from 6 to 300 J and from 0.8 to 6 m/s, respectively. The trolley was instrumented with an opto-electrical device for the measurement of its displacement. This device consisted of a conical needle fixed on the trolley and passing between a light source and a photodiode. The shape of the needle was chosen to ensure that the voltage measured with the diode was proportional to the displacement of the needle. Measurements were done using a numerical oscilloscope and led to force versus time and displacement versus time curves.

Some total impact experiments were performed using an equivalent experimental procedure in which the sample is clamped on the piston of a hydraulic apparatus. This allows measurements at constant velocities up to 10 m/s.

Procedure for the Characterization of the Defied Polymer

The deformed samples were characterized by a description of their shape (diameter, thickness, height), of their microstructure (using optical microscopy on 8-[[micro]meter]-thick slices), of their crystalline organization and orientation (using DSC and X-ray diffraction) and of their resulting mechanical behavior (using a miniaturized tensile machine under an optical microscope).

Thicknesses were determined using a micrometer ([+ or -]1 [[micro]meter]). DSC analysis consisted of scans at a constant heating rate of 10 [degrees] C/min on different zones of the sample using a Perkin-Elmer DSC7 apparatus. The Debye-Scherrer X-ray diffraction technique (flatfilm camera and filtered Cu K[Alpha] radiation) was used to point out crystalline orientation in the material. The incident beam was always perpendicular to the specimen surface.

Tensile tests were performed on specimens taken in different parts of the impacted samples. Because the thickness of the samples was not constant, the specimens were polished to get 100 to 150-[[micro]meter]-thick films of constant thickness. The width of these films was a few millimeters. It was verified that this preparation did not modify the orientation, nor the mechanical behavior. A 1 mm/min cross-head velocity (strain-rate [less than][10.sup.-2] [s.sup.-1]) was chosen to avoid self-heating. Tests were performed at 30 [degrees] C. Tensile specimens were marked with circular dots in order to measure local strains and stresses.


Description of the Impacts

Measurements at high impact energy showed that the energy of rupture for the plaques ranged between 80 and 90 J. Consequently, impacts at energies [less than]80 J, that is, for our device, for falling heights lower than 49 cm, never led to rupture. In fact, when one considers the shape of the sample [ILLUSTRATION FOR FIGURES 1 TO 7 OMITTED] and the evolution of the force [ILLUSTRATION FOR FIGURES 8 TO 10 OMITTED], the energy scale can be decomposed into six domains:

* For low falling heights (e.g., [less than]4 cm) the deformation of the polymer is almost totally elastic. The sample recovers its initial shape as soon as it has been removed from the clamping device. The force versus time curve is a simple peak [ILLUSTRATION FOR FIGURE 8-1 OMITTED].

* If the impact energy increases (falling heights between 4 and 13 cm, that is, velocities ranging from 0.9 to 1.6 m/s) the sample exhibits a non-recoverable deformation which consists of a little dent whose height does not exceed a few millimeters. This dent is located under the striker. The maximum of the force increases with the falling height [ILLUSTRATION FOR FIGURE 9 OMITTED].

* When the falling height reaches roughly 13 cm (1.6 m/s and 21 J) the dent can be decomposed into two parts [ILLUSTRATION FOR FIGURE 1 OMITTED]: a central zone whose aspect is equivalent to that of the initial polymer, and a transparent ring which surrounds the previous zone.

The force versus time curve exhibits a shoulder [ILLUSTRATION FOR FIGURES 8-2 and 8-3 OMITTED].

* For increasing falling heights (between 20 and 45 cm) the non-transparent zone remains unchanged despite the fact that the total deformation of the sample increases [ILLUSTRATION FOR FIGURES 2 AND 3 OMITTED]. In fact, the resulting shape of the sample consists of a kind of "bubble" with a small pastille at its top. The force vs. time curve reaches a plateau, referenced a in Fig. 8 [ILLUSTRATION FOR FIGURES 8-4 TO 8-6 OMITTED]. The maximum force still increases with impact energy, but in a less important way [ILLUSTRATION FOR FIGURE 9, VELOCITIES BETWEEN 2 AND 3 M/S OMITTED].

* At intermediate heights (from 50 to 66 cm), the central zone progressively disappears [ILLUSTRATION FOR FIGURES 4 AND 5 OMITTED]. At this moment a second plateau is observed in the evolution of the force [ILLUSTRATION FOR FIGURE 8 OMITTED]. Depending on the sample, a rupture may be observed under the striker [ILLUSTRATION FOR FIGURE 6 OMITTED], as the final step of this stage. In that case the rupture occurs during the second plateau [ILLUSTRATION FOR FIGURES 10-1 AND 2 OMITTED].

* At high impact energy (in our case, energy [greater than]90 J) rupture appears before the central zone disappears and during the first plateau [ILLUSTRATION FOR FIGURE 10-3 OMITTED]. Rupture is located in a zone which surrounds the striker [ILLUSTRATION FOR FIGURE 7 OMITTED]. Then, the general shape of the sample and the force curve do not significantly depend on impact energy and, provided that impact energy is sufficient, results obtained using the falling mass device are equivalent to those from the hydraulic apparatus [ILLUSTRATION FOR FIGURE 9, VELOCITIES HIGHER THAN 3 M/S OMITTED].

Consequently, at this stage the deformation of the polymer seems to result from the initiation and the propagation of a neck. This local deformation is initiated in a zone which is not located under the striker. This can be explained, as already suggested or observed, by the friction conditions between the polymer and the striker (7, 9, 10, 14). In the case of a frictionless contact, promoted by lubricating the striker, the deformation, and as a consequence the rupture, occur under the striker. This was verified in our case.

When impact energy increases, the maximum force significantly increases till a significant necking occurs. Then, it changes little. During a high energy impact, one may suspect that necking is initiated just before the maximum of the peak, as already suggested (13). As the neck develops, a decrease in the force down to a plateau is recorded. The top of the sample remains almost unchanged in a first step. Then, there is a competition between enlargement of the neck toward the top and rupture in the bottom of the bubble, which would lead to the second plateau and a crack under the striker. Generally, at high impact velocity, rupture occurs before this second step.

Characterization of the Deformed Polymer: Location of the Deformation

The deformation of the polymer was quantified using the characteristic dimensions shown in Fig. 11:c is the diameter of the undeformed zone, e is the height of the bubble, a is the diameter of the deformed zone once the bubble has been laid flat, d and b are the minimum and the maximum diameters of the bubble.

The dimensions of the bubble are reported in Fig. 12. As already mentioned, c remains unchanged at low impact energy and decreases down to zero for energies ranging from 50 to 90 J. e and a monotonically increase with increasing impact velocity. At energies higher than 134 J, samples exhibit the same features as in Fig. 7, that is, rupture occurs at the bottom of the "bubble" and c remains equal to 2-3 mm. d and b are always lower than the diameter of the striker. For example, for falling heights higher than 30 cm (for which the entire hemispherical head of the striker penetrates the sample), d and b are equal to 18 and 19 mm, respectively, whereas the striker has a diameter of 20 mm. This indicates an important elastic recovery at the end of the impact. This also promotes some waving at the bottom of the sample [ILLUSTRATION FOR FIGURES 4 AND 5 OMITTED].

The thickness of the samples varies from a location to another [ILLUSTRATION FOR FIGURE 13 OMITTED] and depends on the impact energy. Two domains are observed. For low impact energy [ILLUSTRATION FOR 4 AND 3 IN FIGURE 13 OMITTED] the thickness of the top of the sample only slightly decreases, whereas the thickness of the surrounding polymer decreases to a large extent. The deformation is then concentrated around the striker as already mentioned for polyethylene (5). For higher impact energy [ILLUSTRATION FOR 2 AND 1 IN FIGURE 13 OMITTED] the pastille progressively disappears. At this moment the thickness of the surrounding polymer remains almost unchanged. Consequently, what has been called "propagation of the neck" in the previous part of the paper consists in fact of two steps:

* First an initiation step in which the thickness of the zone around the pastille decreases down to a minimum value (200-300 [[micro]meter]). During this step, which corresponds to the first plateau in Figs. 8 and 10, the deformed volume remains of the same order of magnitude.

* Second, a propagation stage in which the central zone disappears. This second step can be considered as the actual propagation step and corresponds to the second plateau in Figs. 8 and 10.

Characterization of the Deformed Polyamide 6: Evolution of the Polymer

Samples 2 to 5 mm wide were taken either in the longitudinal (parallel to the impact direction) or in the transverse direction [ILLUSTRATION FOR CIRCULAR SAMPLES IN FIGURE 11 OMITTED]. Longitudinal samples were analyzed on numerous locations (every 2 mm) to determine their crystalline orientation. Both longitudinal and circular (at different heights) samples were tested in tension under a microscope.

The observation of the microstructure, at the end of the first stage of the necking, using optical microscopy on thin sections (8 [[micro]meter]) reveals that the spherulites are no more visible in the transparent zone [ILLUSTRATION FOR FIGURE 14 OMITTED]. The undeformed zone is identical to the initial polymer. The transition zone (i.e., the neck shoulder) is narrow.

Despite important variations of the traces from a sample to another, DSC analysis reveals some evolution in the crystalline structure of the polymer:

* The deformed polymer exhibits a slightly higher enthalpy of fusion (5% to 10%), that is, seemingly, a higher crystallinity ratio or a higher perfection of the crystalline phase;

* Initial polymer most often exhibits only one melting peak (maximum at 228 [degrees] C) and rarely a second peak at 220 [degrees] C.

* This second peak is more often observed in the deformed polymer.

In parallel, the initial plaque mostly gives X-ray diffraction patterns with only one reflection (diffraction angle of 10.59 [degrees]). A second reflection, less visible, seems to exist for an angle of 5.49 [degrees] and is more often visible in the deformed polymer. Polyamide 6 is known to crystallize in different metastable monoclinic structures whose parameters vary between two stable phases (15): the [Alpha]-phase whose parameters are 9.56, 17.24 and 8.01 [Angstrom], respectively, and whose angle between a and c axis is 67.5 [degrees], and the [Gamma]-phase whose parameters are 9.33, 16.88 and 4.78 [Angstrom], respectively, and whose angle between a and c axis is 59 [degrees]. In these unit cells, b corresponds to the chain axis. [Alpha]-phase appears during low cooling-rate crystallization whereas [Gamma]-phase is related to crystallization during flow (e.g., in melt-spinning) (16). The latter phase has a melting temperature lower than the former (214 and 222 [degrees] C, respectively (15)). It would be tempting to conclude that during impact polyamide develops a new crystalline phase closer to its [Gamma]-phase. Unfortunately, our measurements do not allow a precise conclusion on that point. Nevertheless, it can be stated that the crystalline organization of the polymer seems to be different in the deformed zones of the impacted sample and in the initial polymer.

More convincing are the measurements concerning crystalline orientation and mechanical behavior. The initial polymer is isotropic and ductile, whereas the behavior of the impacted polymer depends on the location in the sample: at the top, in the middle or at the bottom of the bubble.

The Top of the Sample:

The pastille, when it exists, seemingly remains isotropic, as indicated by continuous Debye-Scherrer rings [ILLUSTRATION FOR FIGURE 15-1 OMITTED] and in agreement with the microscopic observations [ILLUSTRATION FOR FIGURE 14 OMITTED]. Conversely, the polymer immediately surrounding this pastille is highly oriented [ILLUSTRATION FOR FIGURE 15-2 OMITTED]. This orientation is mainly uniaxial, parallel to the impact, even if a small amount of biaxial orientation can be observed: the reflection at 10.59 [degrees] is mainly equatorial and the one at 5.49 [degrees] mainly meridian. The mechanical behavior of the pastille remains ductile [ILLUSTRATION FOR FIGURE 16-2 OMITTED] but is different from that of the initial polymer [ILLUSTRATION FOR FIGURE 16-1 OMITTED]. The Young modulus in the pastille is 150 MPa compared to 120 MPa for the plaque. The longitudinal Young modulus in the oriented polymer, close to the pastille, ranges from 300 to 500 MPa. Plastic flow is less important in that zone [ILLUSTRATION FOR FIGURE 16-4 OMITTED]. This feature is quite representative of oriented polymers.

When the pastille disappears [ILLUSTRATION FOR FIGURE 5 OMITTED], and if the sample does not break, the uniaxial orientation of the upper oriented zone disappears. The top of the sample becomes biaxially oriented, as indicated by continuous diffraction rings. Logically, the behavior of this zone is between that of the pastille and that of the uniaxially oriented polymer [ILLUSTRATION FOR FIGURE 6-3 OMITTED]. When rupture occurs at the top of the sample [ILLUSTRATION FOR FIGURE 6 OMITTED], the crystalline orientation around the rupture (e.g., at the top of the bubble) is once again uniaxial, but perpendicular to the previous one.

Medium Part of the Sample (Half Height):

As diffraction experiments are done further from the top of the sample (every 2 mm), the degree of uniaxial orientation first decreases. The medium part of the bubble seems to exhibit a more or less pronounced biaxial orientation [ILLUSTRATION FOR FIGURE 15-5 OMITTED]. Circular samples taken in that zone have a Young modulus of 350 MPa and actually behave as an oriented polyamide [ILLUSTRATION FOR FIGURE 17-3 OMITTED], with a high level of transverse orientation.

Lower Part of the Sample:

Uniaxial orientation increases again and is maximum at the bottom of the sample [ILLUSTRATION FOR FIGURE 15-9 OMITTED], whenever the rupture occurs or not, and wherever it occurs. This part, when tested in circular samples cut perpendicular to the orientation, exhibits a ductile behavior which is intermediate between those of the initial polymer and the pastille (i.e., between curves 16-1 and 16-2).

Consequently, the entire zone of the polymer concerned by the impact is modified. The modifications may consist of morphological and crystalline changes, and of crystalline (or molecular) orientation. The deformation is inhomogeneous either from a geometrical (thicknesses) or from a microstructural point of view (orientation). A drastic evolution can be observed on narrow zones. Owing to strain-hardening processes the mechanical behavior of the polymer also varies and is strongly inhomogeneous. In a general manner, Young's modulus increases and ductility decreases. The rate of variation depends on the location with respect to the top of the striker. Although quantitative relationships remain difficult to establish these modifications in the mechanical behavior can be related to microstructural evolution, and especially to crystalline orientation. Part of the sample remains isotropic, part is almost uniaxially oriented, part is multiaxially oriented. Figure 18 summarizes all these observations.

This complex evolution is the trace of a complex loading path. The deep microstructural modification induces important anisotropy effects.


When impacts are performed on PA 6 without lubrication the deformation results from the initiation and the enlargement of a neck.

The first step of the initiation process occurs around the striker just before the force recorded on the striker reaches its maximum. This initiation goes on with the reduction of the thickness of the neck, whereas the force decreases and even reaches a plateau. The zone located under the striker remains almost undeformed and unchanged in that step. Four zones can be described in the sample:

* The zone under the striker whose deformation is not sufficient to induce neither any significant change in the spherulitic morphology nor any apparent crystalline orientation. Nevertheless, this region does not behave as the initial polymer: it is somewhat more rigid, showing that some strain-hardening process has already occurred.

* The zone immediately surrounding the previous one. The polymer is mainly uniaxially oriented parallel to the impact. It no more exhibits any observable semi-crystalline microstructure and behaves as an oriented polymer (higher modulus and lower ductility).

* The zone located at the bottom of the sample, which is also almost unaxially oriented parallel to the impact and which does not contain spherulites. Its behavior is close to that of the former zone.

* As the striker moves forward, the polymer that has been first uniaxially deformed is gradually stretched in the transverse direction owing to the shape of the striker. Consequently, the polymer is more and more biaxially oriented in the mid-height of the deformed sample, and as a consequence, its behavior in that zone is significantly different.

In a second step, the pastille zone under the striker may disappear, whereas the thickness of the neck does not vary. Owing to the symmetry of the loading, this zone becomes biaxially oriented and its behavior is modified. This step is revealed by a second plateau on the force versus time curve. When rupture occurs in that zone, the orientation measured after rupture is parallel to the rupture direction. A possible interpretation could be the initiation of a secondary necking which propagates in a random direction. Then cracking occurs perpendicular to this second necking.

If impact energy is sufficient, as in actual impacts, rupture occurs in the lower part of the sample before the step of propagation occurs. Then rupture is observed around the striker in the bottom oriented zone.

All these observations confirm that the microstructural evolution of the polymer during impact is important and varies from a location to another. Resulting mechanical properties are different. To model impact in a correct way, an accurate constitutive equation should be able to predict and to take into account such different evolution and orientation. Obviously, simple uniaxial deformation, and consequently orientation, is not able to correctly represent the polymer since in an impact situation the sample orientation may be uniaxial, biaxial or nonexisting. This is the main reason why the models for impact are generally not satisfactory in the case of ductile polymers.


The authors want to gratefully acknowledge Vincent Morisseau and Gabriel Monge for their help in the experimental work.


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Title Annotation:French Research on Structural Properties of Polymers
Author:Billon, N.; Haudin, J.M.
Publication:Polymer Engineering and Science
Date:Oct 1, 1997
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