Polymers under mechanical stress: deformation of the nanostructure of isotactic polypropylene revealed by scanning force microscopy.
The plastic deformation of semicrystalline polymers is of great importance to many technological processes such as uniaxial drawing of films, biaxial stretching and solid state forming (1). Because of the structural complexity of those polymers, the understanding of their macroscopic mechanical behavior requires to study the plastic deformation mechanisms that occur at different levels: spherulites, stacks of crystalline lamellae, individual lamellae, molecular chain, and so on (2, 3). Much experimental work has been devoted to the mechanical behavior of semicrystalline polymers and many elementary deformation mechanisms have been proposed (4-8). However, in many cases, evidence of the mechanisms has been obtained by indirect experimental techniques such as X-ray diffraction and birefringence (9-14). In contrast with the case of single crystals and melt cast films (15-17), only few direct observations have been made on bulk deformed samples: by transmission electron microscopy (TEM) (18, 19), by scanning electron microscopy (SEM) (20) or by scanning force microscopy (SFM), also called atomic force microscopy (21, 22). Up to now, most of the scanning force microscopy experiments have been performed on ultradrawn oriented polymers to visualize and to characterize the nanofibrils (21, 22); little attention has been given to clarify the mechanisms of plasticity in the crystalline part of bulk semicrystalline polymers.
In this paper, our aim is to show that scanning force microscopy is a powerful tool to study the local plastic mechanisms which occur in bulk semicrystalline polymers at moderate plastic strains. SFM is used here to investigate at the lamellar level the evolution of the crystalline morphology of intruded isotactic polypropylene (iPP) samples under moderate applied shear stresses. Indeed, since the glass transition temperature [T.sub.g] of iPP is low, its amorphous phase deforms elastically at room temperature and its yield stress is mainly governed by the plasticity of the crystalline phase.
Among synthetic high polymers, iPP is known to have a most complex crystalline structure. Numerous experimental works have been devoted to the study of iPP crystalline architecture and have shown the existence of different levels in the structural hierarchy. All these studies have been performed on isothermally crystallized samples (single crystals, thin films and bulk samples as well). Two major crystalline forms, namely the monoclinic [Alpha] one and the hexagonal [Beta] one have been identified (23, 24). [Alpha]-spherulites exhibit a cross-hatch type lamellar branching where radial and tangential lamellae are oriented nearly orthogonally (25, 26). Depending on the isothermal temperature of crystallization, [T.sub.c], two types of [Alpha]-spherulites can be distinguished: [[Alpha].sub.I]-spherulites nucleate at "low" [T.sub.c] ([T.sub.c] [less than or equal to] 136 [degrees] C for thin films (25)) and contain a great fraction of tangential lamellae which does not vary noticeably with [T.sub.c], [[Alpha].sub.-II]-spherulites are formed at higher [T.sub.c][[T.sub.c] [greater than or equal to] 137 [degrees] C for thin films (25)] and reveal a progressive reduction in the proportion of tangential lamellae with increasing temperature. In contrast, [Beta]-spherulites do not exhibit tangential lamellae and consist of broad, locally stacked radial lamellae, just as in the spherulites of other polymers. Norton et al. (25) have shown that, depending on the [T.sub.c]-range, two kinds of [Beta]spherulites could be encountered: [[Beta].sub.II]-ones whose lamellae look like extended sheet structures in all lateral directions and [[Beta].sub.IV]-ones whose lamellae are twisted periodically along the growth direction and give rise to some concentric banding in polarizing optical microscopy (24).
For iPP, the cross-hatched lamellar pattern of [Alpha]spherulites (25) suggests strongly that the plastic deformation mechanisms should be different than for polyethylene. Recently, the plastic deformation modes have been investigated at the spherulitic level by SEM in bulk intruded iPP samples under uniaxial tension and simple shear (20). Under tensile loading, it has been shown that the [Alpha]-spherulites exhibit brittle behavior whereas the [Beta]-ones deform plastically up to strains of about [Epsilon] = 0.25; the brittle behavior of the [Alpha]-spherulites is characterized by the occurrence of cavitation at early stages of strain near the spherulites boundaries and along the equatorial region perpendicular to the tensile axis. Under shear stress, the [Alpha]-phase cavitation disappears and both [Alpha] and [Beta] spherulites are capable of undergoing plastic deformation (up to strains of about [Gamma] = 1); however, [Alpha]spherulites deform [approximately]25% less than the whole specimen (for [Gamma] = 1) and that lack of plastic deformation is compensated by the [Beta]-spherulites. Unfortunately, because of the low resolution of SEM, the authors were not able to investigate the deformation mechanisms at the lamellar level inside both [Alpha] and [Beta] spherulites. Our attempt here is to correlate the global behavior of each type of spherulite with the plastic events occurring inside them at the lamellar level by the aid of scanning force microscopy. Intruded i-PP samples have been subjected here to moderate strains by simple shear.
Indeed, because the deformation in polymers in the shear mode is macroscopically homogeneous, without any discontinuities like necking observed in tensile experiments, the development of deformation textures, the lamellar morphology changes and the plastic processes can be followed step by step.
The iPP was manufactured by Appryl (France). It is characterized by a broad molecular weight distribution as assessed by gel permeation chromatography with [M.sub.n] = 75,940 and [M.sub.w] = 262,000. The iPP pellets were processed by intrusion in a thick mold designed for producing parallelepipedic iPP plates (300 x 200 x 15 mm). The intrusion process consists in slowly extruding the melt (whose initial temperature is 230 [degrees] C) in a mold kept at 30 [degrees] C and continuing the mold feeding under the extruding pressure (6 MPa) during the cooling sequence (cooling time = 240 s). The cooling kinetics is slow enough to ensure a reproducible semicrystalline structure. Such intruded plates have no detectable orientation and possess much less chain degradation than injected ones. Because of the temperature gradient between the two external surfaces, the crystallization kinetics varies across the thickness of the plates and Aboulfaraj et al. (27) have shown that the spherulitic structure is thickness-dependent. In the mid-thickness of the intruded plates, both [Alpha] and [Beta] spherulites are present, the fraction of [Beta]-spherulites is [approximately]60 vol % and their mean size is 120 [[micro]meter] while their fraction is 0% at the external surfaces of the plates. The iPP samples were cut out of the central area of the plates; they have a parallelepipedic shape with a longitudinal notch 5 mm deep and 4 mm wide [ILLUSTRATION FOR FIGURE 1 OMITTED]. L, d and e are the dimensions of the calibrated part and their values are given in Fig. 1. This geometry favors the localization of the plastic deformation within the predefined plane at the root of the notch. The fiat surface has been polished with several emery papers and finally with a very fine alumina powder (0.05 mm) until no residual scratch was visible. Simple shear tests were performed in an Instron machine at room temperature with a strain rate of 5 x [10.sup.-4][s.sup.-1]. A typical experimental stress-strain curve [Sigma]([Gamma]) is shown in Fig. 2. To enhance the presence of the crystalline phase, the samples were immersed for some time at room temperature in an acid solution (1.3 wt% KMn[O.sub.4], 32.9 wt% [H.sub.3]P[O.sub.4] and 65.8 wt% [H.sub.2]S[O.sub.4]), which etches preferentially the amorphous phase (28). Then, the samples were rinsed in a dilute [H.sub.2]S[O.sub.4] solution, in hydrogen peroxyde, in distillated water and in acetone; and, finally, dried in a vacuum oven. To compare our data with the previous SEM ones (20), that acid attack procedure has been used before the shear test for 40 min. The chosen attack time corresponds to the minimum time necessary to get a good, homogeneous lamellar resolution by SFM; i.e. the whole amorphous phase is removed from the surface by the acid attack, the surface is thus crystalline, and the crystalline lamellae can be imaged through the topographis contrast.
SFM experiments were conducted on a Nanoscope III scanning force microscope (Digital Instruments). Measurements were carried out in the "Tapping Mode." In that mode, the cantilever is forced to oscillate at a frequency (331 KHz) close to its resonance frequency with an adjustable amplitude. Mean value of the repulsive normal force was [approximately]0.1 nN. The surface of the sample can be imaged in two different ways. "Height" images are obtained by using the feedback loop, which maintains the amplitude at a constant value; height measurements are deduced from the vertical displacement of the piezoelectric scanner and the image reflects the topography of the sample surface. "Amplitude" images are obtained when the feedback loop is not connected; the amplitude can vary and the image corresponds to its variation. Since we were interested in the surface topology, all the images shown here are "Height" images and they have been filtered through the "Planefit" procedure.
Figure 3 is an optical micrograph of the surface of an undeformed sample. It shows the crystalline structure at the spherulitic level: the black spherulite is a [Beta] while the white spherulites are [Alpha] (27).
SFM observations of undeformed intruded iPP samples have been performed for two main reasons: i) to compare the crystalline structure of intruded samples (intrusion is not an isothermally procedure) with one isothermally crystallized; ii) to test the SFM technique by comparing observations with previous ones made by TEM (25).
The lamellar structure of both [Alpha] and [Beta] spherulites has been investigated by SFM. Figure 4 is a (560 x 560 nm) SFM image of the center of a [Alpha]-spherulite: the core is an intimately woven array with large amounts of "cross-hatched" lamellae and its structure looks "nodular" rather than lamellar. In Fig. 4, the two main directions of the cross-hatched texture lie approximately along the two diagonals of the SFM image. When progressing outward from the centre (along the radial growth direction: [a.sup.*] axis) (29), there are increasing radial lamellae although large areas are infilled with the tangential component as shown in Fig. 5. Figure 5 is a typical (2 x 2 [[micro]meter]) SFM image obtained in an area located at about the mid-radius of a [Alpha]spherulite; the core of the spherulite lies down the right bottom corner of the image. The radial lamellae are oriented along the diagonal that goes from the right bottom corner to the left top; the tangential lamellae are clearly visible in several fiat areas of the SFM image. At the periphery of the spherulite, high resolution SFM observations (images about 600 x 600 nm in size) reveal a lamellar structure similar to that of Fig. 4; a slight decrease of the tangential component is however noticeable. Nevertheless, the proportion of tangential lamellae is high from the center of the spherulite to its periphery. The mean thicknesses of the radial and tangential lamellae have been measured from (2 x 2 [[micro]meter]) SFM images: their values are [approximately]22 nm and 19 nm, as reported (31). The mean angle between the radial and tangential lamellae has been measured from (2 x 2 [[micro]meter]) SFM images; its value is [approximately]85 [degrees] in agreement with the value of 80 [degrees] 40 [minutes] predicted through crystallographic considerations (26). In intruded iPP specimens, the structure of [Alpha]-spherulites is thus similar to that of [Alpha]-spherulites in isothermally crystallized samples; i.e. a network of radial lamellae with intertwined tangential lamellae as mentioned by Norton and Keller (25). Both the amount of tangential lamellae and the lamellar thicknesses observed here in the [Alpha]-spherulites of intruded iPP correspond approximately to the morphology of the [[Alpha].sub.I]-phase of iPP thin films crystallized isothermally at [approximately]130 [degrees] C.
The structure of the [Beta]-spherulites as revealed by SFM in intruded iPP samples has been reported (31), and, here we will just recall briefly the main characteristic features. In contrast with the [Alpha]-spherulites, the [Beta]-spherulites do not contain tangential lamellae, but radial ones only. In their equatorial plane, the [Beta]-spherulites exhibit a "core" made up of stacked edge-on lamellae surrounded by two lobes of stacked flat-on lamellae (31). Because of the acid attack, the roughness of the polymer sample is high enough to make impossible the achievement of large scale SFM images (about the spherulite size) with a homogeneous lamellar resolution. However, smaller scale SFM images (3 [[micro]meter] x 3 [[micro]meter] in size) obtained along the same radial direction of a spherulite show that, apart from the core, most of the lamellae tend to twist along the radial growth direction and present a broad spread of orientations varying from b-axis (edge-on lamellae) to c-molecular axis (flat-on ones) profiles; flat-on and edge-on lamellae are seen alternately with all the states in between. The mean measured lamellar thickness is [approximately]25 nm (31). Owing to the limited field of view of the SFM, it is not possible to conclude if the observed twisting is periodical or not; however, refering to previous works (24, 25, 30), it seems that the observed [Beta]-spherulites are of type [[Beta].sub.IV]. As mentioned previously (31), numerous hexagonal etch terraces indicating possible screw dislocation growth mechanisms are visible as well as rounded or faceted flat-on lamellae. In isothermally crystallized thin films, the [[Beta].sub.IV] phase occurs at about 130 [degrees] C and, consequently, the morphology of both the [Alpha] and [Beta] spherulites observed in the mid-plane of the intruded iPP plates corresponds to that which is found in thin films isothermally crystallized at [approximately] 130 [degrees] C.
On the other hand, it is clear that the SFM observations confirm previous TEM experiments on iPP (25) and are thus reliable. We can expect that scanning force microscopy will be a good tool to visualize the local plastic mechanisms at the lamellar level in iPP.
Samples Deformed Under Shear Stress
To analyze the first steps of plastic deformation, all the intruded iPP samples have been deformed to a permanent shear strain [Gamma] = 0.5, left under load for 15 min and then unloaded.
Let us first consider the case of the [Alpha]-spherulites. Compared with the case of the unsheared samples [ILLUSTRATION FOR FIGURE 5 OMITTED], the resolution at the lamellar level is lost in numerous areas of any (2 x 2 [[micro]meter]) SFM image taken in any place of a spherulite. High resolution SFM scans (about 600 x 600 nm in size) have been performed in the regions where the lamellar resolution could be reached; those images reveal a lamellar structure similar to that of Fig. 4. Basically, at [Gamma] = 0.5, there are no noticeable changes in the lamellar structure; no plastic events (lamellar kinkings, shear bands, and so on) have been detected inside the spherulites. The 'blurred" aspect of the SFM images of the [Alpha]-spherulites is still not well understood. To clarify that point, two sets of experiments are in progress. The first set consists in carrying out the acid attack procedure after the shear test and not before it; the aim is to check if that "blurred" aspect of the SFM images could be related to the absence of the amorphous phase at the external surface. In the second set of experiments, the iPP samples are sheared at higher permanent shear strains in order to test if the loss of lamellar resolution observed at [Gamma] = 0.5 is some kind of transient effect. The corresponding data will be published in a next future.
The only salient feature is the presence of very few "cracks" located generally along the interfaces between the [Alpha]-spherulites [ILLUSTRATION FOR FIGURE 6 OMITTED]; those "cracks" appear preferentially at the interfaces that lie parallel to the shear direction and sometimes penetrate into spherulites along the radial lamellae when they are also parallel to the shear axis. Such "cracks" were not visible in the undeformed sample.
That behavior can be easily explained by the interlocking effect of the radial lamellae by the tangential ones inside each spherulite. At that level of shear stress, in areas where there are numerous [Alpha]-spherulites, the sample can deform plastically in fragile zones only, i.e., along the spherulites boundaries that contain a lot of defects.
Let us now focus on the case of the [Beta]-spherulites. Figures 7 and 8 illustrate typical (4 x 4 [[micro]meter]) SFM images obtained along the principal tensile strain axis (oriented at [approximately]40 [degrees] of the shear direction for [Gamma] = 0.5 as shown in [ILLUSTRATION FOR FIGURE 1 OMITTED]): collective nanocracks (10-100 nm in width) perpendicular to the tensile strain axis are observed going through stacks of edge-on lamellae [ILLUSTRATION FOR FIGURE 7 OMITTED] and through stacks of flat-on lamellae as well [ILLUSTRATION FOR FIGURE 8 OMITTED]. In the case of edge-on lamellae, the nanocracks propagate along the molecular axis. Along the principal compressive strain axis (defined in [ILLUSTRATION FOR FIGURE 1 OMITTED]), profuse local kinkings of edge-on lamellae are visible [ILLUSTRATION FOR FIGURE 9 OMITTED]. Only few nanocracks parallel to the molecular axis are detectable here and there in individual edge-on lamellae parallel to the shear direction [ILLUSTRATION FOR FIGURE 10 OMITTED]. Incidentally, cracks parallel to the shear direction are seen along the edge of stacks of flat-on lamellae [ILLUSTRATION FOR FIGURE 11 OMITTED].
In a given spherulite, the plastic glide inside the lamellae is favored in the areas where the chains are inclined at 0 [degrees] and 90 [degrees] with respect to the shear direction; nevertheless, within the SFM resolution, no trace of chain slips (which are the precursors of shear bands) has been detected inside the lamellae parallel or orthogonal to the shear direction.
It is important to mention that, at [Gamma] = 0.5, no particular preferred lamellar orientation has been detected in sheared [Beta]-spherulites, and that, globally, except for the above-mentioned local events, the lamellar structure approximately resembles that for the undeformed samples (31).
Since the lamellar morphology of the [Beta]-spherulites of iPP is close to that of the spherulites of polyethylene (32), which similarly contain radial lamellae only, the plastic deformation mechanisms occurring inside the spherulites should be similar in both polymers; it is thus interesting to compare the above results with earlier experiments performed on polyethylene.
In linear high-density polyethylene (HDPE) deformed plastically either by uniaxial compression (18) or by simple shear (19), it has been pointed out that, for moderate permanent strains ([Epsilon] [less than] 0.73 for uniaxial compression and [Gamma] [less than] 0.5 for simple shear), little change in the spherulitic structure is detectable. In both cases, TEM experiments show that no shear bands are present. In HDPE, fine shear bands start to develop only at [Gamma] = 1 in shear tests and at [Epsilon] = 0.82 in compressive ones. In the latter case, TEM experiments reveal clearly the presence of multiple and cooperative kinkings of radial lamellae and the authors point out that the kinking process acts as the precursor of the development of shear bands for higher permanent strains (18). The local kinkings of lamellae we observed along the principal compressive strain axis are consistent with those previous data on HDPE; however, it is clear that, in our experiment, the permanent compressive strain is smaller than that is required to create cooperative lamellar kinkings and to induce plastic shear banding.
The lack of lamellar slip in the [Beta]-spherulites of iPP samples under moderate permanent shear strain ([Gamma] = 0.5) agrees with previous data on HDPE (19); however, our results differ from the data obtained on HDPE by the visualization of nanocracks perpendicular to the principal tensile strain axis. There are two possible explanations: either the existence of the nanocracks at the external surface of the iPP samples is due to the removal of the amorphous elastic phase prior to the shear test or those nanocracks are characteristic of the surface behaviour and do not occur inside the bulk. Indeed, in the first case, at the external surface, only the crystalline phase is subjected to the shear stress; due to the acid attack, the lamellae are no more connected by tie molecules through the amorphous elastic phase and thus the applied stress is poorly transmitted. When the amorphous phase is present, the high molecular mobility, which characterizes it, increases the ability to lower stress concentrations and avoids the propagation of cracks. Here, in the absence of the amorphous phase, the opening of nanocracks could occur in lamellar zones containing defects (more fragile areas) and where the local tensile stress is high as well. The experimental procedure followed by Bartczak et al. (18) was different from ours; the deformed specimens were first stained and fixed in chlorosulfonic acid and ultrathinsections were then cut from the samples for the TEM experiments. In chlorosulphonation, electron-dense atoms are incorporated at lamellar surfaces, the polyethylene is thus converted in a crosslinked polymer with chlorine and sulfur atoms attached to lamellar surfaces (32). The question is whether the staining procedure does modify the local mechanical behavior of the samples. Moreover, it has been shown that the crystallinity falls during chlorosulfonation: the decrease is [approximately]50% for polyethylene crystallized anabarically at 5kbar (32). Since the authors do not discuss that point in their paper (18), it is thus difficult to draw a conclusion from the non-observation of nanocracks by TEM. To make wholly clear the influence of the acid attack before the shear test, SFM experiments are actually in progress on samples first sheared and then acid attacked. Second, in the work of Bartczak et al. (18), the ultrathinsections are cut inside the bulk of the samples and the absence of nanocracks could signify that the bulk behaves in a different way than the surface. To know if the behavior of the external surface reflects the behavior inside the bulk, we plan in the future to cut from the sheared samples parallelepipedic specimens with their fiat surfaces parallel to the external surface of the sample and to observe those fiat internal surfaces by SFM.
For intruded iPP samples deformed under an applied shear stress, we have shown that scanning force microscopy is able to image directly the influence of the macroscopic applied shear stress on the spherulitic structure at the lamellar level. Only one level of permanent shear strain ([Gamma] = 0.5) has been investigated; the corresponding preliminary data are encouraging and promising. Compared to the previous SEM data (20), scanning force microscopy has revealed nanoscopic features that were imperceptible by scanning electron microscopy. In the next future, we plan to investigate higher permanent shear strains ([Gamma] [greater than or equal to] 1) to detect the formation of shear bands. The point is to check if shear bands emerge at the external surfaces of sheared semicrystalline samples. Up to now, as shown in the literature, profuse shear banding has been detected at the external surfaces only in amorphous glassy polymers. By investigating a wide range of shear and tensile strains, SFM experiments should make it possible to visualize the different plastic deformation mechanisms active at different strain levels at the free surfaces and in the bulk: the second step would be to compare them with theoretical models (1).
The authors are very grateful to R. Seguela (LSPES, USTLille) for several illuminating discussions. This research was supported in part by the Contrat de Plan Etat-Region 1994-98, by the FEDER (CEE program) and by the Ministere de l'Enseignement Superieur et de la Recherche Francais (MESR).
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|Title Annotation:||French Research on Structural Properties of Polymers|
|Author:||Castelein, G.; Coulon, G.; G'sell, C.|
|Publication:||Polymer Engineering and Science|
|Date:||Oct 1, 1997|
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