Recyclability Of A Layered Silicate-thermoplastic Olefin Elastomer Nano Composites

Polymer Degradation and Stability 91 (2006) 2396e2407 www.elsevier.com/locate/polydegstab

Recyclability of a layered silicateethermoplastic olefin elastomer nanocomposite M.R. Thompson*, K.K. Yeung MMRI/CAPPA-D, Department of Chemical Engineering, McMaster University, 1280 Main Street West, Hamilton, Ontario, Canada L8S 4L7 Received 5 December 2005; received in revised form 27 March 2006; accepted 30 March 2006 Available online 12 May 2006

Abstract A multiple-pass study was undertaken with a layered silicateethermoplastic olefin elastomer (TPO) nanocomposite to study the impact of processing history on the properties of the material. A set of 10 passes were completed through a co-rotating intermeshing twin-screw extruder with samples collected to monitor changes in the composite. The microstructure of the nanocomposite was characterized using TEM, XRD, FTIR, steady and complex shear rheology, and mechanical testing. With progressive passes through the extruder, the TPO nanocomposite experienced both delamination of the organoclay as well as thermo-oxidative degradation. The onset and extent of degradation were found to be unaffected by the presence of the organoclay species in the polymer, though, inclusion of a maleated compatibilizer led to increased chain scission. The generated carbonyl groups along the polymer chain as a result of oxidation were speculated to have a significant effect on the developing percolating network of clay within the material and on the final rheological properties of the composite. Despite the occurrence of degradation in the nanocomposite during recycling, its rheological and mechanical properties remained significantly higher than those of the unfilled resin. Ó 2006 Elsevier Ltd. All rights reserved. Keywords: Thermo-oxidative degradation; Nanocomposite; Recyclability; Organoclay

1. Introduction One polymer family of growing interest to the automotive industry are thermoplastic olefin elastomers (TPOs) which are based on polypropylene and an elastomeric resin such as ethyleneepropyleneediene terpolymer (EPDM), ethylenee propylene rubber (EPR), or an ethyleneea-olefin metallocene copolymer. Both compounded and reactor grades exist for TPO materials. Variation of the rubber content with respect to the polypropylene as well as the addition of fillers and process additives gives end-users the ability to target use of the TPO to different interior and exterior applications throughout the vehicle. The use of fillers in a TPO, such as talc, increases the mechanical properties and scratch resistance of the molded

* Corresponding author. Tel.: þ1 905 525 9140x23213; fax: þ1 905 521 1350. E-mail address: [email protected] (M.R. Thompson). 0141-3910/$ - see front matter Ó 2006 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymdegradstab.2006.03.013

part, yet additionally, increases weight which is typically considered a negative factor in automotive design. Polymerlayered silicate (PLS) nanocomposites are an emerging class of polymers requiring considerably less filler content to achieve similar or improved mechanical properties [1e7]. The layered silicate most commonly chosen for these composites is montmorillonite, a clay mineral comprised of layered silicate sheets. The surface chemistry of the montmorillonite is altered with an organic surfactant, to increase its hydrophobicity and allow for improved interactions with polymer species. The critical issue to the manufacture of nanocomposites is the in situ generation of individual nanometer-scale silicate sheets (platelets) within the polymer matrix through a process of intercalation and delamination. Effectively, the original organoclay particles must be broken down to primary particles, and polymer chains be allowed to enter and expand the clay galleries leading to ultimate dissociation of the individual silicate sheets [1e3,8,9]. The nanometer-scale dimension of the exfoliated filler avoids the presence of stress

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concentrators in the matrix which are regions prone to mechanical failure with conventional fillers. Despite the advantageous properties of these new polymers, they remain difficult to synthesize due to the suspected importance of diffusionlimited intercalation of polymer chains into the confined regions within the layered silicate structure [10e12]. It is the unique nature by which PLS nanocomposites are generated that also complicates issues of recycling of this material whether its source comes from post-consumer plastics or production scrap. Knowledge of microstructural changes arising from re-processing of a PLS nanocomposite is important for evaluating the economical and potential environmental impact for this relatively new class of material. This paper will limit its scope of investigation to that of recycling production scrap by examining the morphological response of a TPO nanocomposite subjected to multiple deformation/thermal histories. Prior studies [3,5,6,13e20] regarding the morphology and properties of PLS nanocomposites derived from a polyolefin matrix, namely polyethylene and polypropylene, have highlighted the difficulties with non-polar polymers producing exfoliated structures. The mechanism of intercalation necessary for exfoliation to occur is a diffusion-dependent process reliant on residence time and the interactions between polymer and clay to overcome the subsequent entropic penalty resulting from confinement of the intercalant [10e12,21]. To overcome the unfavorable approach of non-polar polymers into the galleries of the polar organophilic-coated montmorillonite, a polymeric compatibilizer must be introduced in a small quantity sufficient to improve interactions with the silicate surface yet remain miscible with the matrix [3,17]. The combination of an organoclay and polymer compatibilizer has been sufficient to generate partial intercalated/exfoliated structures for polyolefin-based nanocomposites yet the mechanism of intercalation remains hindered. The repercussions of hindered intercalation among non-polar resins such as TPO mean that under subsequent processing histories, as would be expected from recycling, there remains the high likelihood of further intercalation/exfoliation. Effectively, repeated recycling of such materials can be likened to annealing of the nanocomposite in the presence of shear. Mechanical and rheological properties would be expected to shift accordingly with changes in the extent of intercalation/exfoliation. In addition to changes in the microstructure of the montmorillonite, recycling is expected to bring about changes in the molecular structure of the matrix as a result of degradation. The only other known work examining recyclability of a nanocomposite was conducted by Lew et al. [2], though in their case the chosen matrix was a polar polyamide-12, which being a polar macromolecule more readily intercalates the layered silicate than a non-polar polymer. In their work, delamination continued for seven passes through a single-screw extruder with simultaneous crosslinking resulting from thermo-oxidative degradation. This work attempts to examine the impact of recycling on intercalation/exfoliation and degradation for a TPO nanocomposite. A multiple-pass method, commonly used to study degradation in processing equipment, is uniquely applied in this paper to follow the relatively slow progression of intercalation

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of a non-polar resin into the clay structure. It is the intent of the study to reveal the interactions of the polymer and organoclay during recycling. Use of a TPO in this study was significant since this resin is one of the earliest examples of a commercial nanocomposite; however, the fact that it is a blend of uncertain composition presents some complications in the present analysis. Particularly, some questions regarding the participating polymer species in the occurring degradation of this work will not be fully addressed due to the proprietary nature of the material. 2. Experimental 2.1. Materials The chosen matrix was a reactor-grade TPO (Basell Hifax CA387A). The resin had a melt index of 19 g/10 min (230
C/2.16 kg ASTM D1238). For propriety reasons, no attempt was made to determine the composition of the material besides a DSC scan to establish the melting temperature of the material. The DSC confirmed that the major component was an isotactic polypropylene and through our experience with comparable resins we know that the elastomer phase will likely make up less than 25 wt% of the alloy. The organoclay used in the synthesis of our nanocomposite was Cloisite 15A (Southern Clay Products), a natural montmorillonite modified with a hydrogenated tallow substituted quaternary ammonium salt. XRD analysis of the clay filler noted a basal spacing of ˚ confirming values provided by the supplier. d001 ¼ 32.57 A The selected compatibilizer (MAdPP) was a functionalized polypropylene, namely Polybond 3200 supplied from Crompton Corporation, containing 1 wt% grafted maleic anhydride. 2.2. Sample preparation Melt intercalation of the polymer nanocomposite was conducted on a Leistritz 40 L/D ZSE-27 HP co-rotating twin-screw extruder configured for strand pelletization. The trials required two different extrusion configurations. The initial configuration, used only for the first pass, involved compounding the organoclay into the polymer. A side stuffer attached to the extruder metered the clay into the molten polymer at a distance of 18 L/D from the feed opening. Two screw designs were used in this work, a high-work mixing screw and a moderate-work mixing screw, both shown in Fig. 1. The philosophy behind the design of the high-work screw was to generate substantial dispersive mixing in order to break-up agglomerates and to uniformly disperse the clay and compatibilizer within the TPO matrix. This screw was only used for the first pass of our polymer through the extruder. The moderate-work screw design was used for the subsequent nine passes through the extruder and focused more heavily on distributive mixing with consideration for both the work needed for melt intercalation and the need to minimize degradation. Previously reported studies [1,3,8,14] have indicated that melt intercalation favors distributive mixing with a moderate level of shear.

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M.R. Thompson, K.K. Yeung / Polymer Degradation and Stability 91 (2006) 2396e2407 Table 1 TPO formulations Sample name

TPO resin (wt%)

MAePP (wt%)

Organoclay (wt%)

VTPO TPOP TPOC

100 91 88

0 9 9

0 0 3

Number of passes through the extruder.

coupled phenomena. The extruder was operated with a flat barrel temperature profile (180
C for the first pass and 200
C for subsequent passes) and a constant screw speed (200 rpm for the first pass and 75 rpm for subsequent passes). Colorimetric tracer analysis revealed that the mean residence time was 32 s for the first pass and 95 s for each subsequent passes through the extruder for these operating conditions. 2.3. Characterization

Fig. 1. Schematic diagrams of the (a) moderate work and (b) high-work screw designs used in the 40 L/D, 27 mm co-rotating intermeshing twin-screw extruder.

Three resin systems were compounded in this study according to the formulations listed in Table 1. The TPO/MAePP blend and the neat TPO were compounded along with the nanocomposite to provide a means of analyzing the aspects of intercalation versus degradation which are otherwise

Complex viscosity (h*) and the storage modulus (G0 ) were determined in the range 0.1e100 s1 for samples from each pass through the extruder using an ARES (TA Instruments) parallel plate rheometer. A constant strain of 5% was chosen for the measurements based on the determined linear viscoelastic region for our materials from strain sweep testing. The tests on the parallel plate rheometer were conducted at 210
C. A ROSAND dual-bore capillary rheometer was used to determine the shear viscosity of the samples for the shear-thinning region between 100 and 8000 s1 at a temperature of 210
C. Rabinowitsch and Bagley corrections were used for all capillary measured data. Film samples were prepared using a hot press at 180
C under 350 bar pressure for 2 min. These films were characterized at 2 cm1 resolution in a Nicolet 520 FT-IR spectrometer to observe changes in the different functional group concentrations based on number of passes. To evaluate the extent of local intercalation that occurred after passing through the extruder, samples of the TPO nanocomposite were studied using a transmission electron microscope (TEM) and X-ray diffraction (XRD). Direct observation of the layered silicate structure in the polymer matrix was possible using a Philips CM12 TEM operating at an accelerating voltage of 120 keV for 45K, 60K and 125K magnification. For TEM micrographs taken at the 106 magnification, samples were analyzed on a JEOL 2101F field emissions electron microscope. Specimens were prepared by using a diamond knife to microtome thin sections (nominally 70-nm thick) perpendicular to the flow direction from the pelletized samples. Analysis by XRD was performed on a Bruker D8 diffractometer using Cu Ka radiation. The scanning rate was 2
/min in wide angle mode for the 2q range from 0 to 10
. Mechanical properties were determined for the different samples on an INSTRON 3366 testing machine. All test specimens were prepared on a 55-ton Arburg injection molding machine. The flexural modulus was measured by a 3-point bending test according to ASTM D-740 using a strain rate of 0.10 mm/min and cross-head speed of 1.43 mm/min. A

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17:1 support span ratio was used for the bending tests. The tensile properties (i.e. Young’s modulus, yield strength, and elongation at break) were determined according to ASTM D-638 using a 5 mm/min cross-head speed. 3. Results and discussion Typically, researchers studying recycling and the subsequent degradation that occurs within polymers have relied heavily on either rheological measures (e.g. zero-shear viscosity, melt index, storage modulus) or gel permeation chromatography (GPC) to determine changes in the molecular structure. However, this presents a challenge in determining the level of degradation that occurred in a nanocomposite as a result of re-processing or normal service life. Molecular weight determination by GPC becomes infeasible due to the presence of the nano-scale clay particulate which will either affect the inlet filters of the instrument or more likely pass through the filters and proceed to plug the columns and interfere with downstream detectors. Rheological analysis is normally the more sensitive measure of changes in the molecular structure of a polymer compared to chromatographic techniques [22]; however, no previous studies exist to demonstrate whether on-going intercalation can be distinguished from degradation in the case of a PLS nanocomposite. A necessary aspect of this study was, therefore, to assess the capability of rheology to study degradation in this new class of polymers. For clarity of discussion, we will use the terms agglomerate, primary particle, crystallites (tactoids) and platelets (sheets) in the same manner as previously defined [11]. 3.1. TEM and XRD characterization The high extent of shear and extensional deformation characteristic of the screw design used for the first pass required the polymer composite to experience the shortest possible residence time in order to minimize degradation; however, it was recognized that high shear and short residence time are not conducive to intercalation of the organoclay. The purpose of the first pass through the extruder was, therefore, to breakup the majority of all clay agglomerates prior to subsequent passes. Observations made by TEM, as exemplified by the micrograph in Fig. 2(a), of samples collected after the first pass through the extruder suggested that the morphology of the material could be described as possessing an incomplete exfoliated structure (exfoliated platelets with intercalated tactoids remaining). The micrograph showed dark lines indicative of tactoids with thicknesses ranging from 9 to 20 nm as well as grey silhouettes indicative of individual platelets dispersed in the matrix. The tactoids were often observed to be distorted rather than rigid flat planar structures in the matrix. High magnification of the tactoids (refer to the micrograph of a sample from Pass 1 in Fig. 3) showed that expansion of the basal spacing began at the edges suggesting that the intercalant (i.e. matrix polymer or compatibilizer) acted as a wedge disrupting the crystallite structure. Within the interior of the tactoid structures, the layer order appeared to be undisturbed.

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There was no evidence within the TEM micrographs that any clay agglomerates remained after the first processing pass by the TPO nanocomposite through the extruder. The high level of shear of the initial screw design had been sufficient to break-up the clay into primary particles and further into separated tactoids. The subsequent micrographs presented in Fig. 2(b)e(d) show the progression of intercalation and delamination within the composite over time under lower shear conditions than experienced by the material during the first pass through the extruder. The moderate-work screw was designed to maximize residence time and distributive mixing in order to provide optimal conditions for intercalation [1,2,8]. The micrographs show that an increasing number of platelets were present with each successive analyzed pass and that conversely, the remaining tactoids displayed less silicate sheets in their stack structure as well as an increasing frequency of basal expansion localized at their edges (more fraying similar to that shown in Fig. 3). The change in thickness of the tactoids shown by TEM was quantified by image analysis software (SigmaScan) using the value of the minor axis for the determined area of each tactoid. The variation in thickness for the clay particulate within each nanocomposite sample over the different passes through the extruder is shown in Fig. 4 as a set of histograms. During earlier recycling passes the tactoids showed progressive reduction in their thickness through continuous delamination. However, the similarity of the distributions between the seventh and tenth passes indicated that further delamination beyond the seventh pass was not achieved. Further analysis of the structure of the tactoids shown in Fig. 2 made note of the constant gallery spacing regardless of the number of passes through the extruder. While XRD measurements of the basal spacing can be less accurate at determining the actual gallery size due to defects in the stacking structure [10,11], the data confirmed the TEM observations that the d-spacing of the organoclay filler was effectively constant over all 10 passes regardless of the extent of residence time and deformation experienced. A constant d-spacing for the organoclay during delamination was similarly noted by Lele et al. [18] with polypropylene as the matrix resin. It appeared to the authors that the tactoids were being split into thinner crystallites generally near the mid-plane. The tactoid dimensions did not appear to change significantly between the seventh and tenth passes, suggesting that any further delamination of the clay would be minor. Shear-induced splitting of the tactoids (observed within the TEM micrographs) appeared to be the significant mechanism for dispersion of the filler within our non-polar matrix since this phenomenon would maintain the interlayer spacing while diminishing the thickness of the individual tactoid. It is unclear why the interior layer order of the clay was not notably changed by intercalation; however, based on clay colloid chemistry it is reasonable to speculate that the negative dipole at the anhydride oxygen precluded its adsorption into the negative interlayer surfaces of the clay and so the anhydride group would remain located at the clay edges which are the more favorable location for associations with polyanions [23,24]. It

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Fig. 2. TEM micrographs at 60K magnification of the TPO nanocomposite for (a) Pass 1, (b) Pass 4, (c) Pass 7, and (d) Pass 10.

appears, however, that the intercalating compatibilizer or matrix resin may only need to expand the edge of a tactoid to initiate shear-induced splitting. Certainly, among the tactoids seen under TEM that were being sheared, fraying of the edges was always observed. Dennis et al. [1] proposed that for

exfoliation to occur within systems like ours where marginal compatibility existed between the clay and polymer, both shear and chemical compatibilization are necessary. In the case of our TPO nanocomposite, the compatibilization strategy appeared to be inadequate, despite the high weight fraction of additive already employed, and the morphology seemed incapable of progressing to a fully exfoliated structure.

3.2. Steady and complex viscosity

Fig. 3. TEM micrographs at 106 magnification of the TPO nanocomposite after the first pass through the twin-screw extruder.

The complex and steady shear viscosity of both unfilled TPOs (i.e. neat and PPeMA blend) and the TPO nanocomposite are plotted in Fig. 5 with a notable discontinuity between the two sets of measurements. The discontinuity highlights the failure of these materials to satisfy the Coxe Merz rule. Failure of the rule implies that the microstructure of these materials is sensitive to the mode of deformation applied; a phenomenon correspondingly noted by other researchers for reinforced and elastomer-modified polymers [12,18,19,25,26]. At the lowest shear rate in the figure, we find that the nanocomposite exhibited a substantially larger viscosity (for the first pass, 120% over the TPO blend and

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Fig. 4. Histograms of particle thickness for the organoclay obtained by image analysis of TEM micrographs of the TPO nanocomposite for (a) Pass 1, (b) Pass 4, (c) Pass 7, and (d) Pass 10. Number-average (tn ) and weight-average (tw ) thickness values for the distributions are included.

144% over the neat TPO) owing to restrictions in long chain segmental motion within the newly developed percolating network structure of tactoids [19]. As seen in the figure, this shear rate lies in the Newtonian region for the neat TPO and TPO/PPeMA blend; however, with the inclusion of the clay the Newtonian plateau was shifted below the measurement range of our rheometer. In the shear-thinning region (particularly for shear rates larger than 103 s1),

Fig. 5. Flow curves of the three polymer systems determined from parallel plate and capillary rheometers.

differentiation based on viscosity of the nanocomposite from the other two polymers was much more difficult owing to shear-induced disruption of the clay network and the subsequent alignment of tactoids and platelets in the direction of flow [18,26]. Between the first and tenth passes (representing a cumulative residence time of 15 min in the extruder), the TPO/PPeMA blend experienced the greatest decrease in viscosity followed by the neat TPO and finally the nanocomposite. In the case of the unfilled TPOs, the reduction in viscosity was readily attributed to thermo-oxidative degradation. With polypropylene making up the bulk of our material, the decrease in viscosity was explained by the occurrence of b-scission reactions. The ethylene comonomer present in the elastomer phase may have led to some extent of crosslinking, though its effects (if present) were not observable in comparison to the influence of scission on viscosity. For the nanocomposite, the effects of degradation on viscosity were less prominent, either compensated to some degree due to progressive tactoid shearing and exfoliation or delayed due to the clay inhibiting oxidization. In order to determine the role of the organoclay in degradation of the nanocomposite, it is necessary to attempt to distinguish the effects of the two correlated phenomena, namely delamination and degradation, within the rheological data. A key question for this recyclability study lies with whether the clay influences thermooxidative degradation; it has been suggested in the literature that degradation appears to a higher extent in PLS

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nanocomposites compared to a neat resin, due to interference by the clay on the performance of antioxidants [7,13,27]. 3.3. Extent of degradation First, the two unfilled TPO systems were examined to determine the extent of degradation in the absence of the clay. The pass-by-pass changes in the molecular structure of these two materials were approximated by their zero-shear viscosity (h0) (refer to Table 2). The zero-shear viscosity for the TPO nanocomposite was not included in Table 2 since a Newtonian plateau was not reached within the shear range of our rheometer. The neat TPO and TPO/MAePP blend both showed a decrease in their zero-shear viscosity over the 10 passes through the extruder; the apparent onset of degradation being noted for the TPO blend during the third pass through the extruder. The extent of degradation (u) was estimated from this data according to:  3:71 DM M 0  M h u¼ 0 ¼ ¼ 1  00 0 M M h0

ð1Þ

where M0 and h00 is the molecular weight and zero shear viscosity of the polymer, respectively at t ¼ 0 and M corresponds to the molecular weight at some time during processing. This measure of degradation uses the approximate power-law relationship between weight-average molecular weight and zeroshear viscosity [28], but it is considered to be only a basic indicator since it neglects the polydisperse nature of the polymer in the calculations. Using Eq. (1), we estimate that the overall extent of degradation after 10 passes for the TPO blend was more than twice than that for the neat TPO, i.e. u ¼ 0.26 versus u ¼ 0.12, respectively. It has been noted [29] that composites containing the maleated compatibilizer consumed antioxidants more readily and demonstrated degradation sooner in oxidative studies compared to non-compatibilized systems. In addition, it is known from studies examining the free radical grafting of maleic anhydride onto polypropylene [30] that grafted maleic anhydride may act as a chain transfer agent, further promoting degradation reactions. We surmise that the increased susceptibility of samples containing the maleated compatibilizer arises from the labile hydrogen on the attached Table 2 Zero-shear viscositya Extrusion pass

Neat TPO (Pa s)

TPO/MAePP blend (Pa s)

1 2 3 4 5 6 7 8 9 10

3136  218 e e e e e e e e 2001  48

3476  379 3670  321 2132  48 2285  254 1925  58 1956  70 1553  399 1478  32 1396  193 1204  58

a

Estimated by cross model using complex viscosity data.

succinyl anhydride groups along the polymer backbone. Likely, the same susceptibility is experienced by the TPO nanocomposite which also includes the compatibilizer but its influence cannot be readily observed due to the coupled interactions between degradation and clay delamination. To investigate the evolving microstructure caused by the coupled interactions of degradation and clay delamination in the nanocomposite, the storage modulus was analyzed to study changes in chain mobility which may be elucidated from variation in shear rate. 3.4. Change in storage modulus with recycling Fig. 6 compares the storage modulus (G0 ) for the three polymer systems over a frequency range between u ¼ 0.1 and u ¼ 100 rad s1. The plot in Fig. 6(a) shows the modulus decreasing with each of the successive pass through the extruder for both neat TPO and the TPO blend. The general reduction in G0 indicated greater chain mobility, which has already been attributed to chain degradation, predominantly via b-scission. No observable evidence in the curves was present to suggest that the elastomeric phase was affected by recycling; close examination of the low frequency region did not show the increase in G0 noted by other authors [31] which they attributed to crosslinking and aggregation of the elastomer phase in their impact-modified polypropylene. Similarly, the tensile and flexural moduli discussed in a later section of this paper did not show evidence of significant crosslinking within these materials. Fig. 6(b) compares the storage modulus of the polymer material with and without clay in the matrix. At low frequencies (u < 1 rad s1), the storage modulus of the nanocomposite was significantly larger in value compared to the TPO/MAe PP blend for all samples and exhibited a slope close to 0.5 in the terminal region by the fourth pass. The change in the slope of the curve within the terminal region of the graph and the corresponding increase in modulus have been identified by others [12,16,18,19,26] as corresponding to shortrange fillerefiller interactions. It is our belief that crosslinking similarly had no significant contribution to the storage modulus of the nanocomposite as found with the TPO/PPeMA blend. The only possible case by which the extent of crosslinking could be conceivably greater for the nanocomposite than the TPO blend would arise if the filler participated in degradation as a radical transfer agent or initiator but research results in the literature indicate that organoclays do not participate in radical reactions with the polymer [32]. Analysis in the subsequent sections of this paper would appear to corroborate that the organoclay does not participate in thermo-oxidation reactions. At high u, the storage modulus decreased as the number of passes through the extruder increased, though the change was minor for the nanocomposite in comparison to the TPO blend. The low frequency and high frequency behaviors of these materials in the plots of Fig. 6 showed distinct differences which can be related to different aspects of chain motion within the microstructure during oscillatory testing. It is postulated that

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Fig. 6. Storage modulus of the (a) TPO/MAePP blend compared to the neat TPO resin; and (b) TPO nanocomposite compared to the TPO/MAePP blend.

the effects of the clay network and degradation on the rheology of the polymer may be partially separated through their frequency-dominated behavior. At this point we are simply looking for a tool to address our need to determine whether the nanoclay additive is actively participating in thermo-oxidation reaction.

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a filled polymer exhibiting a percolating network under oscillatory deformation have been similarly noted by other authors [33,34], albeit for larger scale particles. In view of the rheological work of Wu and Zheng [33] which examined the sensitivity of the storage modulus in detecting variations in the percolating network of a carbon black filled polyethylene at different frequencies, a frequency u ¼ 100 rad s1 should reasonably minimize the effects of clay in our measurements. The effects of b-scission on G0 , as already shown in Fig. 6, are observable across the full frequency range of the rheometer resulting in a downward shift of the curve. However, its effects will only be observable within the measurements when the effects of the clay network are effectively diminished because of the small magnitude of change brought about by scission. Using the highest and lowest frequencies capable of our parallel plate rheometer, we can attempt to separately discuss apparent changes to the clay network structure (i.e. delamination) and the polymer microstructure (i.e. degradation) for the nanocomposite based on the frequency specific storage modulus. Fig. 7 shows G0 at u ¼ 0.1 rad s1 (to represent the low frequency region) and u ¼ 100 rad s1 (to represent the high frequency region) plotted versus the number of passes through the twin-screw extruder for the TPO/PPeMA blend and TPO nanocomposite. The standard error of G0 for the blend was 4.4%, while for nanocomposite the error was slightly less at 3.6%. Looking at the high frequency (u ¼ 100 rad s 1 ) G0 data in Fig. 7, it was apparent that the trends were similar between the two different polymer systems. For the first two passes through the extruder, the G0 did not change and then for each subsequent passes, the storage modulus decreased. The difference in magnitude between the two polymers indicated that the influence of the percolating network was never fully negated at this frequency else they should match, though the fact that the trend for the nanocomposite does not bear any resemblance to its curve at low frequency suggests that the clay had very little affect. Linear regression applied through the data between Passes 2 and 10 indicated that the rate of decrease for the storage modulus at this

3.5. Degradation and clay delamination in the nanocomposite It is generally stated that under oscillatory shear, large-scale motion of polymer chains predominates at low frequencies, yet become largely restricted at higher frequencies with only local segmental motion being possible. Following this line of thought, the restrictive nature of the percolating clay network on the movement of the polymer matrix can be expected to significantly influence the storage modulus at low frequencies, producing a highly elastic response in the material. However, the influence of this network structure will become increasingly difficult to detect due to shorter relaxation times as the frequency increases. These viscoelastic features of

Fig. 7. Variation of the storage modulus at u ¼ 0.1 and 100 rad s1 with respect to number of extrusion passes by the TPO blend (TPOP) and TPO nanocomposite (TPOC) (lines included for clarity).

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frequency was 1607 Pa/pass and 1609 Pa/pass for the TPO/ PPeMA blend and nanocomposite, respectively. The similarity in the change of modulus at high frequency suggests that degradation progressed at nearly equal rates for both the materials. This would mean that the clay did not have any significant effect on the stability of the polymer. It is most likely that any perceived susceptibility of a nanocomposite to thermo-oxidative degradation is attributed to the compatibilizer rather than the clay. The low frequency (u ¼ 0.1 rad s1) G0 data in Fig. 7 for the TPO/PPeMA blend only displayed a steady decrease in value with increasing number of passes, mirroring the results at high frequency with regards to the influence of chain scission on the polymer structure. For the nanocomposite, the effect of the clay was seen as having a dominant effect on the G0 at this frequency, and the influence of chain scission was considered to be indiscernible based on arguments given later in the discussion. The G0 (u ¼ 0.1) trend for the nanocomposite displayed a steady increase in the stiffness of the material up to the third pass, followed by an abrupt drop till the sixth pass, and then finally exhibited a plateau. The initial increase in G0 for the first three passes through the extruder was attributed to increased fillerefiller interactions as a result of continued delamination during recycling. The XRD and TEM results confirmed that the sizes of tactoids were continually decreased by delamination after the first and fourth passes through the extruder. The increase in the modulus per pass corresponded to decreasing ‘n’ with respect to G0 wun (shown in Fig. 6) as the elasticity of the material became increasingly independent of frequency (below 10 rad s1). The subsequent drop in G0 after the third pass was less apparent as to its cause. The drop in G0 was not likely due to agglomeration of the tactoids since this was not observed by TEM and XRD. Degradation would seem to be the most likely cause for this phenomenon, though perhaps not due to chain scission as might be initially believed. As already argued the modulus contribution derived from the percolating network will be reasonably insensitive to the different relaxation behavior of cleaved chains. This argument is further reinforced by the presence of a plateau in Fig. 7 for the low frequency G0 after the sixth pass. We know from TEM analysis that delamination did not significantly proceed between the seventh and tenth passes, and therefore, in the absence of delamination as a counter-effect to degradation, the decrease in G0 if caused by scission between the third and sixth passes should have continued (or decreased even faster) for the subsequent passes through the extruder. The most plausible explanation for the decrease of the low frequency G0 in Fig. 7 is that fillerefiller interactions of the nanocomposite were reduced due to the increasingly polar nature of the matrix as the nanocomposite was recycled. As thermo-oxidative degradation proceeded with each subsequent passes through the extruder, there would have been an increase in the concentration of oxidation functionalities grafted onto the polymer [35,36] and these carbonyl species could be reasonably expected to interfere with the hydrogen-bonding associations formed within the percolating network. It has been shown in other work [37] that the elasticity of a composite decreases as the compatibility between

the polymer and filler phases is improved. The interference of the degraded polymer species on the network structure is speculated to follow the conceptual associations shown in Fig. 8. Once the polymerefiller interactions increased to a comparable level to those experienced between the adjacent tactoid and platelet particles, the network structure was not further affected and a plateau in the modulus of the material resulted. 3.6. Mechanical properties Mechanical testing was performed under both flexure and tension for the TPO nanocomposite and TPO/MAePP blend to further differentiate the stability of these two materials in regards to degradation, and due to the importance of this data to the automotive industry. The results are summarized in Fig. 9 showing the tensile and flexural moduli, yield strength and elongation at break data for samples of both the materials corresponding to the first, fourth, seventh and tenth passes through the extruder. With the addition of 3 wt% organoclay to the TPO, it can be seen in the figure that after the first pass, the yield strength increased by 11% while the tensile and flexural moduli increased by 30% and 31%, respectively. This significant increase in stiffness at low loadings of filler highlights the importance of this new class of material to the polymer industry, being particularly beneficial to the automotive sector where the continued need for material weight reduction must always be balanced with consideration of structural requirements. The most dramatic change in

Fig. 8. Conceptual representation of the percolating network of tactoids (a) edge-to-edge and edge-to-face associations between the tactoids due to the hydrophilic edges of organophilic clay swollen with a polycation and (b) in the presence of carbonyl functionalized polymer, a competition for association with the clay surface develops.

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Fig. 9. Change in mechanical properties with progressive passes through the extruder.

mechanical behavior that resulted from the inclusion of the clay was at failure, with the percent elongation at break having decreased by 180% after the first pass. All these reported changes in mechanical properties due to the inclusion of organoclay are consistent in magnitude with the data presented by Lee et al. [38] regarding a similar grade TPO. For subsequent passes through the extruder, the TPO materials exhibited changes in mechanical properties (Fig. 9) that are consistent with degradation. For the TPO blend, the flexural modulus indicated a steady decrease in stiffness with continued recycling through the extruder, which ultimately resulted in a reduction by 20% after the tenth pass. The tensile modulus for this material also decreased but to a lesser extent, dropping by only 13% after the tenth pass. The yield strength for the blend dropped slightly (4%) while the percent elongation at break declined by 68%. For the TPO nanocomposite, the moduli (both flexural and tensile) remained relatively constant over the first four passes and then decreased. The fourth pass corresponds to the observable onset of polymerefiller interactions as noted by the rheological testing. Losses in stiffness were considerably less for the nanocomposite in comparison to the TPO blend, decreasing by less than 10% after all 10 passes had been completed. The yield strength of the nanocomposite showed a slight increase with progressive recycling; the small increase could be attributed to on-going delamination of the clay and possibly by a reduction of the elastomeric domain size which has been noted by other researchers [38]. The elongation

behavior of the nanocomposite followed the trend demonstrated by the modulus, dropping in value beyond the fourth pass through the extruder. The mechanical data found for the nanocomposite showed that the value-added benefits of the organoclay addition were retained despite the occurrence of degradation. Interestingly, the mechanical behavior of the reinforced material did not show any increase in stiffness though delamination continued after the first pass; losses in properties were only observed after the fourth pass when greater allowances for chain mobility was observed within the percolating network (similar to our observations in the rheological analysis). 3.7. Infra-red spectroscopic analysis The mid-range infra-red spectra of samples for the TPO blend and TPO nanocomposite are shown in Fig. 10. Spectra corresponding to the first and tenth passes for each polymer system are highlighted in the figure to demonstrate the major changes, if any, to the molecular structure of the matrix as a result of recycling. We surveyed the bending mode vibrational bands between 800 and 900 cm1 (in Fig. 10(a)) for indications of vinylidene generation caused by b-scission while the stretching mode vibrational bands between 1700 and 1850 cm1 (in Fig. 10(b)) were examined for the presence of carbonyl species attributed to oxidation products (i.e. aldehydes, ketones, and carboxylic acid). Despite rheological evidence of thermo-oxidative degradation, it appears that no evidence was found

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Fig. 10. Mid-range transmission FT-IR spectra of the TPO nanocomposite and TPO/MAePP blend, examining two regions (a) 400e1000 cm1 and (b) 1500e2000 cm1.

indicating the presence of oxidation functionalities from these plots. It has been well established in the literature that carbonyl-based functional groups (i.e. aldehydes, ketones, and carboxylic acid) are always generated through the thermooxidative reaction pathway within the range of processing temperature used in this work [35,36]. The most reasonable explanation for the absence of visible functional groups in the infra-red spectra was that their concentration was too low for detection. Either the intensity of vibrations pertaining to other functional groups (particularly the clay and compatibilizer) obscured the peaks being sought or the extent of degradation was low and thus the concentrations of the functionalities were simply below the sensitivity of the instrument. As a point of note, the authors attempted FT-IR analysis of the samples using both film transmission and attenuated total reflectance (ATR) approaches with no differences noted in the findings. 4. Conclusions The recyclability of a current state-of-the-art TPO nanocomposite has been studied in this work and shown to

represent a complex problem with coupled interactions between morphological development of the clay network structure and degradation. The benefit of recycling a non-polar PLS nanocomposite comes from its inability to completely exfoliate the organoclay under normal (single-pass) production. Subsequent processing steps further promote development of a percolating network of tactoids in the matrix thereby minimizing the inevitable impact of thermo-oxidative degradation on rheological and mechanical properties. Inevitably, delamination will eventually cease to alter the tactoid structure of the material; though, a recommended recycling strategy for this class of material might be to introduce both additional antioxidant and compatibilizer for subsequent passes through processing machinery. The onset of degradation appeared to occur following the second pass through the twin-screw extruder and the presence of the organoclay was concluded to have no effect on either the onset or the extent of degradation experienced by the TPO matrix. Rather, it was felt that the maleated compatibilizer increased the extent of degradation within the matrix, likely due to the participation of its grafted functional group as a chain transfer agent. Oxidative modification of the matrix as a result of degradation was felt to have a negative effect on the stiffness of the material derived from the percolating network structure consisting of exfoliated platelets and dispersed tactoids. A decrease and subsequent plateau in the storage modulus between the third and sixth recycling passes through the extruder suggested reorganization of the microstructure of the nanocomposite as the increasingly polar polymer matrix participated in bond associations with the network structure. Ultimately, the study demonstrated that rheological and mechanical properties of the recycled PLS nanocomposite far surpassed the performance of neat material even with only a small addition of organoclay. Acknowledgements The authors wish to thank the AUTO21 Network Centres of Excellences and the Natural Sciences and Engineering Research Council (NSERC) of Canada for their funding for this work. The authors wish to go onto thank Decoma International for the donation of the resin, Drs. Rempel and McManus from the University of Waterloo for their assistance with the FT-IR analysis, and Drs. Botton and Britten at McMaster for their help and technical advice regarding the TEM and XRD analyses. References [1] Dennis HR, Hunter DL, Chung D, Kim S, White JL, Cho JW, et al. Polymer 2001;42:9513e22. [2] Lew CY, Murphy WR, McNally GM. SPE Annu Tech Conf 2004;50:299e308. [3] Wang Y, Chen F-B, Wu K-C. J Appl Polym Sci 2004;93:100e12. [4] Mehta S, Mirabella FM, Rufener K, Bafna A. J Appl Polym Sci 2004;92:928e36. [5] Wang K, Liang S, Du R, Zhang Q, Fu Q. Polymer 2004;45:7953e60.

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