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Rheological study of mixing in molten polymers: 2-mixing of reactive systems
Polymer 45 (2004) 8031–8040 www.elsevier.com/locate/polymer
Rheological study of mixing in molten polymers: 2-mixing of reactive systems Philippe Cassagnaua,*, Franc¸oise Fenouillotb a
Laboratoire des Mate´riaux Plastiques et Biomate´riaux, UMR-CNRS 5627: Inge´nierie des Mate´riaux Polyme`res ISTIL, Universite´ Claude Bernard, Lyon 1, 15 Boulevard Latarjet, 69622 Villeurbanne cedex, France b Laboratoire des Mate´riaux Macromole´culaires, UMR-CNRS 5627: Inge´nierie des Mate´riaux Polyme`res, Institut National des Sciences Applique´es, 17 Avenue Jean Capelle, 69621 Villeurbanne cedex, France Received 12 March 2004; received in revised form 23 August 2004; accepted 15 September 2004 Available online 1 October 2004
Abstract Three reactive polymer systems have been examined with a new mixing device adapted on a classical rheometer in order to investigate reactive mixing situations encountered in polymer blends. After having characterized the bulk polymerization of 3-caprolactone (3-CL), the polymerization of 40 wt% of 3-caprolactone into a copolymer of ethylene and vinyl acetate (EVA) was run into the rheo-mixer. The kinetics of the reaction in dispersed media was observed slightly different from that in bulk since the characteristic time of 3-caprolactone diffusion into EVA is much lower than its time of mixing. On the other hand, it was observed that the molecular weight distribution of the poly(3caprolactone) is broader in dispersed media (Mw Z 15; 000 g moleK1 ; IpZ2.6) than in bulk (Mw Z 17; 000 g moleK1 ; IpZ1.6). A broadening of the molecular weight distribution in dispersed media was pointed out due to the fact that 3-CL monomer is partitioned between the EVA and PCL phases leading to a non-homogeneous concentration of monomer in the reactive phase. The polycondensation of 40 wt% of a epoxy-amine system into a polystyrene matrix was also investigated and the morphology of the resulting material examined. A gradient of structure and conversion was detected in a blend obtained from the assembly of two initially nonreactive layers. The gradient reveals that the amine diffuses faster than the epoxy leading to non-stoichiometry of the reactive functions across the sample. When the blend was polymerized under shear, the kinetic of the reaction remained unchanged regardless the level of shear. However, the morphologies were significantly different, pointing out the importance of the coalescence and droplet deformation phenomena. Spherical droplets were observed at 0.15 sK1, elongated droplets and fibers at 1.5 and 15 sK1. q 2004 Elsevier Ltd. All rights reserved. Keywords: Rheo-mixing; Diffusion; Reaction
1. Introduction The dispersion and/or dissolution at a molecular scale of low viscosity reactants in a polymer of high viscosity is a key for the control of the final macromolecular architecture when one considers a reactive polymer processing technology in the molten state. Such a reactive process in polymer blending can be for example the in situ polymerization of one phase in the presence of one other [1–3] or the grafting of a functional specie onto a polymer chain [4]. Generally
* Corresponding author. Tel.: C33 4 72 44 62 08; fax: C33 4 72 43 12 49 E-mail address: [email protected] (P. Cassagnau). 0032-3861/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2004.09.028
speaking, the time of mixing must be lower than the time of chemical modification and/or polymerization otherwise the local stoichiometry may be different from what is desired, resulting in non-controlled materials. For example, Cassagnau et al. [1] investigated from the torque variation in an internal mixer, the formation of polyethylene/thermoplastic-polyurethane blend (PE/TPU) via in situ polymerization of diols and diisocyanate in a molten polyethylene matrix. They demonstrated that the viscosity ratio, which considerably increases with the extent of the polymerization reaction, appears to be the physical parameter which governs the blend morphology development. In a same way, Meynie´ et al. [5] investigated the polymerization of a thermoset system into a polystyrene matrix. They show that
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the morphologies are strongly dependent, not only on the evolution of the viscosity ratio between the matrix and the dispersed phase, but also on the gel point and crosslinking phenomena of the dispersed phase. Nevertheless, these experiments were conducted in non-well controlled flow as the variation of the torque can be qualitatively correlated only to an apparent viscosity. On the other hand, due to the amount of polymer introduced in such mixer (about 60 g) and the low conductivity of polymer components, self heating phenomenon by viscous dissipation is another non-well controlled parameter. Reaction exothermicity may be another one. Depending on the processing conditions, i.e. apparent shear rate, a overheating of several degree is generally observed that accelerates the reaction. Finally, when mixing a large amount of a low viscosity reactant with a polymer in a internal mixer, the phenomenon of lubrication of the rotor by the liquid is often observed. In this case, the measured torque is apparently zero and no information on the state of mixing is available until the liquid slowly mix and/or react and stops lubricating the rotor. The rheoreactor concept provides an alternate solution to such rheological characterization problems [6]. The rheoreactor is a curved bottom vessel supplied with a helicoı¨dal double ribbon impeller. Nevertheless, such a system adapted to a conventional rheometer cannot be developed for polymeric system of high viscosity due to the torque upper limit authorized by the transducer of the rheometer. Consequently, this mixer-type rheometer can be only devoted to the study of intermediate viscosity liquids such as water/glycerol, salad dressing, polymer solutions or oligomers [7]. To overcome these difficulties, a mixing device called rheomixer has been developed [8] and adapted on a classical rheometer, with the aim of measuring reliable viscosity data on high viscosity fluids during the course of reactive mixing operations. In a previous work we have examined the process of mixing a low viscosity liquid, miscible in a high viscosity polymer [8]. The objective of the present study is to investigate, the influence of the mixing conditions on the polymerization kinetics and the morphology development of molten-reactive dispersed media. For this purpose, we analyze two systems: the polymerization of 3-caprolactone in a copolymer of ethylene and vinyl acetate (EVA) matrix and a the polycondensation of a thermoset (epoxy-amine) in a polystyrene (PS) matrix. For both, we expect to collect new data with the aim of better understanding the phenomena of diffusion, mixing and morphology formation.
2. Experimental 2.1. Experimental set-up The rheo-mixer has been described elsewhere in more details [8]. The rheo-mixer device was designed in order to
carry out the mixing experiments on a Rheometrics Mechanical Spectrometer (RMS800). This mixing device is based on a parallel plate analogy with some vertical obstacles like spikes in the gap for breaking up the streamlines of the fluid. A calibration procedure was used to convert the rotational speed/torque in shear/stress for a quantitative analysis of the torque curves. Most of the tests were conducted in the step shear mode. The viscosity was calculated from the torque values and the apparent applied shear rate according to the calibration constants [8]. Consequently, the mixing curves show the variation of the viscosity as a function of the time of mixing. Some tests were run in dynamic mode in the domain of the linear viscoelasticity. For the latter, a parallel plate geometry with 25 mm diameter was used. 2.2. Materials 2.2.1. Polymers A copolymer of ethylene and vinyl acetate (EVA) with a melt index of 420 was supplied by Atofina. The amount of acetate groups contained in this copolymers is 28 wt%. This copolymer was named EVA28420. Its zero shear viscosity is 360 Pa s at TZ100 8C. The polystyrene (PS) Lacqrene 1450 N was also supplied by Atofina. 2.2.2. Monomers and initiator 3-caprolactone (3-CL) and titanium tetrapropoxide were obtained from Aldrich and Roth Sochiel, respectively. The epoxy was a diglycidyl ether of bisphenol A (DGEBA) with a degree of polyaddition of nZ0.15 supplied by Bakelite. The diamine was 4,4 0 -methylenebis[2,6-diethylaniline] (MDEA) supplied by Lonza. 2.3. Sample preparation Globally, we applied the same experimental procedure described in a previous paper [8]. The particular procedures used for the reactive systems are detailed in the following section. 2.3.1. 3-Caprolactone The initiator was added to the 3-caprolactone just before the experiments in the rheometer. The reactive mixture was transferred, through a 2 ml syringe, to the vessel of the rheomixer preheated to the experimental temperature. The exchange surface between the reactive mixture and the area was very small in comparison with the volume, limiting the loss of monomer by evaporation. 2.3.2. EVA/3-caprolactone The EVA disc was brought into contact with the lower plate of the vessel at room temperature. The oven was then heated to the experimental temperature, TZ102 8C. The reactive liquid (3-CLCinitiator), was transferred in the vessel, through a syringe, over the EVA molten layer. The
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Fig. 1. Bulk polymerization of 3-caprolactone in the rheo-mixer at different shear rates. Rheo-kinetics at TZ102 8C.
polymer and liquid thicknesses were calculated in order to adjust the concentration of poly(3-caprolactone) in EVA at about 40 wt%. The total thickness of the reactive medium (EVAC3-CL) was 1.6 mm. 2.3.3. PS/thermoset The composition of the blend studied was 60 wt% of PS and 40 wt% of epoxy/amine thermoset precursors. In order to avoid a reaction of the monomers during a blending operation, two non-reactive pre-blends were prepared by extrusion. Their composition were calculated to obtain a stoichiometric mixture of epoxy functions with the amine functions of the co-monomer. The thickness of the nonreactive blends discs was adjusted in order to respect the epoxy-amine stoichiometry and the global composition of the blend. Thus: hPS/DGEBAZ0.9 mm and hPS/MDEAZ 0.7 mm. The epoxy rich disk was placed on the lower plate (25 mm parallel plate geometry) with the amine rich disk just above it. The total thickness of the reactive medium was 1.6 mm.
3. Results and discussion 3.1. Bulk polymerization of 3-caprolactone Modeling the change in viscoelastic properties during a reaction of polymerization is of great importance for predicting the flow behavior since flow is strongly dependent on the degree of polymerization. In a previous study [9], the rheological behavior during the bulk polymerization of 3-caprolactone with titanium tetrapropoxide as the initiator was studied directly in the rheometer with parallel plate geometry and in the domain of the linear viscoelasticity. Thus, the low deformation applied could not promote mechanical mixing. The aim of the present work is to study the influence of high deformations and mixing
conditions on the rheo-kinetics of 3-CL bulk polymerization. The length of the polymer chains depends directly on the ratio [M]/[Io]. [M] and [Io] are, respectively, the monomer and initiator concentrations. The polymerization can be classified as living because of the linear variation of the average molecular weights in number and in weight versus the conversion. For the present work [M]/[Io]z600. Fig. 1 shows that a same evolution of the viscosity (rheokinetics) during the polymerization is obtained for shear rates in the range 1.5–150 sK1. The viscosity changes from the viscosity of the 3-CL monomer at the early stage of polymerization to the viscosity of the PCL at the end of the polymerization. However, the description of the viscosity is less accurate for low shear rate, especially at the early stage of polymerization due to the sensitivity of the rheometer sensor. Besides, it can be pointed out that the PCL samples completely polymerized exhibit a Newtonian behavior over the whole range of shear rate investigated here. Furthermore, the molecular weights of the poly (3-caprolactone) samples at the end of the polymerization process in the rheomixer were measured by size exclusion chromatography system (SEC). For the different processing conditions, we observed the same molecular weights of the samples within an experimental error of 10%: Mw Z 17; 000 g moleK1 and Mn Z 10; 500 g moleK1 that is a polymolecularity index of IpZ1.6. Consequently, the fact that the viscosity–kinetic curves depict a same curve at different shear rate is not surprising since the system is thoroughly mixed before its introduction into the mixer. In other words, the polymerization process is carried out in a perfect homogeneous media. Such a result was already reported in bulk polymerization of linear polyurethane in extruder [10,11]. 3.2. Polymerization in dispersed media New developments in reactive polymer blends have
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Fig. 2. Polymerization of 3-caprolactone in the rheo-mixer during mixing with EVA28420 at different shear rates. TZ102 8C. Initial conditions: EVAC3CLZ1.6 mm. EVA/PCL at the end of the mixing and polymerization processes: 60/40 wt%.
increased sharply because reactive compounding may provide viable mechanisms for the in situ elaboration of the desired blend with controlled structure and morphology. However, the complex nature of polymer reactive modification processes in dispersed media requires a fundamental understanding of the mechanisms that govern chemical reaction in polymeric melt phases, as well as a critical assessment of the role of key process parameters on the endusage properties of the blend. The aim of this part is to investigate experimentally the effect of mixing on reaction kinetics and morphology development of reactive polymeric blends produced by polymerization of monomers in dispersed media and in controlled processing conditions (shear rate, temperature). For this purpose, two systems were considered. The first one is the formation of a EVA/PCL blend via in situ polymerization of 3-CL and the second one is a thermoplastic (TP)/thermoset (TS) blends resulting from the in situ polycondensation and of thermoset precursors in a polystyrene phase. 3.2.1. EVA/3-caprolactone We used the same experimental conditions as those used in bulk polymerization (TZ102 8C, [M]/[Io]z600) for the polymerization of the 3-CL in the molten EVA28420 matrix. Fig. 2 shows the variation of the apparent viscosity versus the mixing time (or reaction time) for different shear rates. Remind that the characteristic time of the bulk polymerization of 3-CL is 200 s as observed in Fig. 1. In dispersed media, Fig. 2 shows that the time of polymerization, tp, of 3-CL in EVA48420, is around 300–400 s, that is close to the time observed in bulk. In a previous work [8], we observed for the same processing conditions (temperature and shear rate) that the characteristic time, tm, for mixing 3-CL monomer with EVA28420 was around 3000 s. The comparison of these characteristic times indicates that
the polymerization rate is not really slowed down when the 3-CL polymerizes in presence of the EVA. Thus, the monomer does not completely incorporate to EVA prior to reacting. This is not very surprising since 3-CL and EVA are not miscible in all proportions as was observed visually and may be predicted from their respective solubility parameters (dCLZ21.1 MPa1/2 and dEVAZ18 MPa1/2). Furthermore we demonstrated, that the mixing process is controlled by the diffusion process of 3-caprolactone monomer in molten EVA phase [8]. By analogy with the Deborah number, we can define a dimensionless number as the ratio of the characteristic time of the mixing (tm) to the characteristic time of the polymerization process (tp): Db Z
tm tp
(1)
Here Db z10 confirming that the 3-CL polymerization is mainly performed in a bulk state. Indeed, Fig. 2 shows that the time of mixing/polymerization is in the same order of magnitude whatever the shear rates. The difference in the shape of the curves is due to the fact that the viscosity of EVA28420 and the viscosity of the PCL at the end of the polymerization are in the same order of magnitude (h0,EVA28420Z280 Pa s and h0,P3-Cl Z360 Pa s at TZ _ 1:5 sK1 Þ only a slight 100 8C). At low shear rate ðgZ variation in the viscosity curve is observed. This curve depicts an inflexion point around 100 s. At high shear rate _ 15 and 150 sK1) a slight lubricant effect is observed at (gZ 20 s followed by an increase of the viscosity due to the 3-CL polymerization. Actually, the viscosity ratio considerably changes with the extent of the polymerization, varying from l!10K5 at the early stage of the polymerization and mixing processes to lz1 at the end of these two coupled processes. The
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Fig. 3. Molecular weight distributions of poly(3-caprolactone) synthesized in bulk and in EVA dispersed media (1.5 and 15 sK1).
mixing between EVA and PCL is accelerated due this evolution of the viscosity ratio with the extent of the reaction of polymerization. This discussion is based on the fact that these two polymers are immiscible which has been proved by differential scanning calorimetry study of the present system at the end of the polymerization process. The controlling effect of the viscosity ratio on mixing was noted on other systems as the reactive blending by in situ polymerization of an polyurethane phase in a polyethylene media for instance [1,11]. Fig. 3 shows that the molecular weight distribution of the PCL is broader in dispersed media (Mw Z 15; 000 g moleK1 ;
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IpZ2.6) than in bulk (Mw Z 17; 000 g moleK1 ; IpZ1.6) and that the molecular weight distribution is not dependent on the mixing conditions (shear rate) as observed in Fig. 3. Actually, the PCL and EVA phases are PCL and EVA rich phases containing 3-caprolactone since the monomer is present until the end of the reaction, and can diffuse in EVA. Thus, the broadening of the molecular weight distribution in dispersed media can be explained by the fact that 3-CL monomer is partitioned between the EVA and PCL phases leading to a concentration gradient of the 3-CL in the reactive phase. We should also point out that the initiator of the 3-CL polymerization is also a catalyst of the transesterification reaction so that exchange reactions could be encountered between vinyl acetate groups and 3-CL ester groups. And actually, when the polymerization of 3caprolactone in EVA was carried out in extruder at 160 8C, NMR analysis have shown that 3-caprolactone was grafted onto EVA chains by a transesterification reaction [12]. Nevertheless, the polymerization temperature of the present work is low (TZ1008C) compared to the temperature in extruder (TO160 8C), and the reaction of transesterification can be neglected as was confirmed by NMR analysis. 3.2.2. PS/thermoset This part concerns the polymerization of a epoxy-amine thermoset system (DGEBA–MDEA) into a PS matrix. The
Fig. 4. PS/thermoset (60/40) reactive blending at 177 8C and 60 rpm in a internal mixer. (a) Glass transition of the PS phase (&) glass transition of the epoxy phase (:) and torque (—) as a function of time. Final morphology (b) reaction in an oven and (c) reaction in the mixer.
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Fig. 5. PS/thermoset (60/40). Modulus of the complex shear viscosity at uZ1 rad sK1 and TZ177 8C. Linear viscoelasticity—curve A: initial condition, two layers constituted by PS/DGEBA (hZ0.9 mm) and PS/MDEA (hZ0.7 mm); curve B: initial conditions, one layer constituted by a pre-blend PS/DGEBA– MDEA prepared in a batch mixer at TZ80 8C.
concentration of the thermoset precursors is 40 wt%. This system is quite complex and more details can be found in the work of Meynie´ et al. [5]. In the cited work, the authors considered the evolution of the blend under shear in a batch mixer at TZ177 8C. This study was based on a homogeneous initial blend of epoxy and amine monomers in PS, at least the authors supposed that the concentration in reactants at a micro-scale is at the stoichiometry. Their results confirmed this assumption. At the temperature of the experiments (TZ177 8C), the monomers are soluble into the PS so that the epoxy and amine functions are diluted and the reaction is slowed down compared to a pure DGEBA– MDEA mixture at the same temperature. When the chain length of the growing epoxy-amine species increases, phase separation occurs. A dispersed phase is formed by nucleation and growth. This phase is mainly composed of epoxy-amine oligomers and monomers. This causes an acceleration of the polymerization. Then, the composition of the phases evolves as can be seen from the Tg data, and the epoxy droplets grow in size by diffusion of the monomers and coalescence until the epoxy domains gel and the conversion goes to completion (Fig. 4(a)). When the reaction is run in quiescent conditions, the final particles are spherical (Fig. 4(b)) while under flow, in a internal mixer for instance, they deform and coalesce ending to large, irregular particles (Fig. 4(c)). Note that at the end of the process both the PS-rich phase and the epoxy phase are purified since their Tgs are those of the pure PS and pure DGEBA–MDEA network. In the present work, we aimed to study the influence of the diffusion and of the mixing of epoxy and amine monomers on the morphology development of the reactive blend with PS. A first experiment was run in dynamic regime at uZ 1 rad sK1. The strain amplitude was adjusted from 100% at
the beginning of the test to 1% at the end of the experiment to obtain a measurable torque and to remain within the domain of linear viscoelasticity. In linear viscoelasticity the entities (polymer chains, droplets,.) are deformed around their equilibrium position. In other words, the material structure is not altered during the course of measurement in the domain the linear viscoelasticity [13]. Thus we assume that this test was performed in ‘static’ condition, that is no coalescence was produced by the shear. This assumption is valid since we found that the morphology obtained with the operating conditions described above was similar to the morphology obtained in a oven in the absence of any shear (compare Fig. 6 with Fig. 4(b)). This demonstrates that we can assimilate experiments run in the rheometer in oscillatory mode, to experiments run in truly static conditions. Two different preparations of the sample test were experienced. The first one is a two layers sample constituted of a PS/DGEBA layer and a PS/MDEA layer. Each layer is not reactive unless it is mixed or put in contact with the other one. The second one is a unique layer of a homogeneous PS/DGEBA–MDEA blend prepared at low temperature (80 8C) in order to limit the reaction. We can see the evolution of the complex viscosity of the two different preparations of the samples in Fig. 5 showing the variation of the complex shear viscosity versus reaction time at TZ177 8C. Curves A and B in Fig. 5 exhibit an identical qualitative variation of the viscosity with polymerization time. Note that the viscosity data brings information on the evolution of the polymerization and on the morphology of the blend [14–16]. The phase separation can be well identified for both curves in Fig. 5. When the epoxy phase separates, the PS matrix viscosity rises. It is the first change of the slope of the viscosity variation with time. Later, the gelation is identified by a marked increase of the viscosity.
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Fig. 6. Condition for: 1-two initial layers PS/DGEBA and PS/MDEA (curve A, Fig. 5); (a) sample from PS/DGEBA layer external side, (b) sample from interface between PS/DGEBA and PS/MDEA layers and (c) sample from PS/MDEA layer external side, and 2-one initial layer constituted by one homogeneous PS/DGEBA-MDEA layer (curve B, Fig. 5).
We find the phase separation around 22 min, and the gel point is detected around 37 min. Finally, the rheology of the blend above the gel point of the thermoset phase is essentially governed by the rheology of the PS rich matrix since the thermoset phase has a droplet morphology. The increase in viscosity above the gel results from the migration of amine and epoxy monomers and/or oligomers out of the PS-rich phase toward the epoxy domains. Consequently, the viscosity and the glass transition temperature of the PS matrix increases with the extent of the reaction. The final viscosity in curve A is 10 times lower than the viscosity in reference curve B indicating that the epoxyamine reaction is globally slower, and that the epoxy conversion is not 100% at the end of the process. For the two layers sample, it is clear that the diffusion of the monomers
in the neighbor layer is controlling the process. At the beginning, we note that the polymerization is only slightly slower since the phase separation and the gel are detected, respectively, at 22 and 42 min. This indicates that in this period of time, the diffusion is faster than the polymerization. Latter, the evolution of the viscosity, thus of the epoxy conversion, is much slower. The decrease of the diffusion coefficient of the monomers may be the cause since the global viscosity and glass transition temperature of the phases increase with the conversion of the epoxy, making the polymerization diffusion controlled. The final morphology of PS/DGEBA–MDEA from curve B (one layer) consists of a dispersion of spherical thermoset rich particles of about 3 mm diameter (see Fig. 6.2). A similar structure has been observed by Meynie´ et al. [5] when the polymerization occurred in static conditions.
_ 0:15 sK1 and TZ177 8C. Curve A: initial condition, two layers constituted by PS/DGEBA (hZ Fig. 7. PS/thermoset (60/40) blending in rheo-mixer. gZ 0.9 mm) and PS/MDEA (hZ0.7 mm); curve B: initial conditions, one layer (hZ1.6 mm) constituted by a pre-blend PS/DGEBA–MDEA prepared in a batch mixer at TZ80 8C.
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Fig. 8. Scanning electron microscopy of PS/thermoset blend at the end of _ 0:15 sK1 Þ at TZ177 8C. (a) from the curing test in dynamic condition ðgZ two initial layers PS/DGEBA and PS/MDEA. (a) From initial homogeneous PS/DGEBA–MDEA system: curve B in Fig. 7. (b) From two initial layers PS/DGEBA and PS/MDEA: curve A in Fig. 7.
Morphology analysis in Fig. 6 shows that a different result is observed for the droplet morphology of the initial two layers PS/DGEBA–PS/MDEA. For the analysis of these morphologies, it must be kept in mind that the epoxy and amine reactants are initially separated so that the reaction may happen only if the monomers diffuse to the adjacent layer. We can assume that diffusion occurs essentially in one direction (axial direction of the sample). The morphology at the external side of the initial PS/DGEBA layer consist in hard thermoset droplet with an average diameter about 5 mm (Fig. 6.1(a)). In the region, initially at the interface between the PS/MDEA and PS/DGEBA layers (Fig. 6.1(b)), droplets with smaller size about 2 mm are observed. The morphology at the external the side of the initial PS/MDEA layers (Fig. 6.1(c)) consists in sub-micron particles with an average diameter about 0.2 mm. This structure is typical of a late nucleation showing that the epoxy concentration was low in
this zone of the sample so that the epoxy droplets did not have time enough to grow in size. From the above remarks, we can deduce that the diffusion of the amine in the PS/ epoxy layer is much faster than the diffusion of the epoxy in the PS/amine layer. Consequently, the DGEBA and MDEA reactants are not at the stoichiometric balance over the sample thickness. Moreover, it can be assumed from this gradient of morphology, a gradient in phase compositions leading to particles with different glass transition temperature. On the other hand, the non-stoichiometric condition implies that some DGEBA and MDEA monomers and/or oligomers remains in the PS phase (in the PS/amine layer) at the end of the polymerization process. As a consequence the conversion is not complete and this explains why the viscosity in curve A is 10 times lower than the viscosity in reference curve B after 80 min of reaction. Fig. 7 shows the variation of the absolute complex viscosity versus polymerization/mixing time at a low shear _ 0:15 sK1 Þ: Curves A and B were obtained from the rate ðgZ same initial conditions as early described in Fig. 5. For an initial homogeneous reactive blend, Fig. 8(a) shows a homogeneous morphology as expected, with droplets diameter of about 10 mm. The droplets remain spherical under weak dynamic conditions but their diameter is larger than that under static conditions: 3 mm (Fig. 6.2). The balance between the mechanism of coalescence and break up is shifted to the coalescence process and leads to larger particles. If we come back to Fig. 7, before the gel, the same evolution of the viscosity is observed compared with Fig. 5. However, after the gel, the viscosity of the two layers sample (curve A) reaches the same value than in curve B at the end of the polymerization (hz105 Pa s). According to our previous discussion, this result means that above a certain viscosity, even a very low shear rate is efficient enough to induce a mechanical mixing that decreases the distance of diffusion of the epoxy to such an extend that the
Fig. 9. PS/thermoset blending in rheo-mixer at different shear rate, TZ177 8C. Initial conditions: one layer (hZ1.6 mm) constituted by a pre-blend PS/DGEBA–MDEA prepared in a batch mixer at TZ80 8C.
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Fig. 10. PS/thermoset blending in rheo-mixer at different shear rate, TZ177 8C. Initial conditions: two layers constituted by PS/DGEBA (hZ0.9 mm) and PS/MDEA (hZ0.7 mm).
polymerization kinetics becomes almost identical to that of the homogeneous sample. Nevertheless, the viscosity is a macro-scale measure as the whole volume of the sample is solicited in the experiment. Morphological analysis in Fig. 8(b) shows that the morphology, resulting from mixing of initial two layers, is not homogeneous at a micro-scale level. From this figure two distinct zones can be observed, the first one is constituted by thermoset droplets with 10 mm diameter and the second one is constituted by smaller droplets with a size lower than 1 mm. The size of these zones is approximately 100 mm. From these results, we conclude that the rheo-mixer working at a low shear rate is efficient to get a macro-scale mixing (scale about 100 mm) for the present system. Figs. 9 and 10 show the variation at high shear rates _ 1:5 and 15 sK1) of the absolute complex viscosity (gZ versus the time of mixing/polymerization for initial homogeneous PS/DGEBA–MDEA blend and initial two PS/DGEBA and PS/MDEA layers, respectively. Just after the phase separation process the viscosity decreases
Fig. 11. Morphology of PS/thermoset blend at the end of the curing test in dynamic condition from initial homogeneous PS/(DGEBA–MDEA) system. For morphology at 0.15 sK1 see Fig. 8(a); (a) 1.5 sK1 and (b) 15 sK1.
depending on the shear intensity. At the gelation process _ 15 sK1 it rapidly the viscosity suddenly increases. For gZ drops few minutes later. Actually the sample becomes _ 0:15 sK1 the unstuck from the mixer plate. Above gZ microscopy analysis (Figs. 11 and 12) shows a rough morphology with a fibrous structure. Some thermoset ellipsoids of 50 mm length can be observed in these figures. The coalescence and deformation of the droplets has produced an elongated structure. Finally, viscosity curves and morphology pictures do not show any noticeable difference between both initial conditions (homogeneous reactive system and non-homogeneous two layers) so that the mixing efficiency of the rheo-mixer is rather good for two initial liquids with a viscosity ratio close to one.
4. Conclusion The reactive mixing process of three different systems has been investigated with the help of a new mixing device. Reactive systems such as bulk 3-caprolactone polymerization and dispersed media polymerization of linear thermoplastic (poly 3-caprolactone), and thermoset (epoxy/amine) were investigated. We showed that the polymerization in EVA molten polymer is slightly slowed down at high shear rates by a dilution effect resulting from the mixing of 3caprolactone in EVA phase. It was observed that the molecular weight distribution of the poly(3-caprolactone) is broader in dispersed media (Mw Z 15; 000 g moleK1 ; IpZ 2.6) than in bulk (Mw Z 17; 000 g moleK1 ; IpZ1.6). This broadening of the molecular weight distribution in dispersed media was pointed out due to the fact that 3-CL monomer is partitioned between the EVA and PCL phases leading to a non-homogeneous concentration of monomer in the reactive phase.
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the Scanning Electron Microscopy work and Sophie Dumas, PhD student, for her help in the polymerization of the 3-caprolactone. The authors also acknowledge one of the reviewers for their very useful comments.
References
Fig. 12. Morphology of PS/thermoset blend at the end of the curing test in dynamic condition from two layers constituted by PS/DGEBA (hZ 0.9 mm) and PS/MDEA (hZ0.7 mm). For morphology at 0.15 sK1 see Fig. 8(b); (a) 1.5 sK1 and (b) 15 sK1.
For the PS/thermoset system, we have used two initially non-reactive layers of PS/DGEBA and PS/MDEA. The experiment run in oscillatory mode, that is in the absence of mechanical mixing allowed to show that the apparent polymerization rate was progressively slowed down in the two layer system. This was attributed to decrease of the diffusion coefficient of the monomers caused by the variation of the glass transition temperature of the phases. Besides, the amine was found to diffuse faster than the epoxy leading to a gradient of morphology and conversion in the two layer sample. The morphology of the blends polymerized under increasing levels of shear depicted spherical thermoset droplets only at very low shear rate. Above 0.15 sK1, epoxy elongated droplets and fibers were produced.
Acknowledgements The authors gratefully acknowledge Pierre Alcouffe for
[1] Cassagnau P, Nietsch T, Bert M, Michel A. Polymer 1998;40:131–8. [2] Cartier H, Hu GH. Polymer 2001;42:8807–16. [3] Wollny A, Nitz H, Faulhammer H, Hoogen N, Mu¨lhaupt R. J Appl Polym Sci 2003;90:344–51. [4] Dassin S, Dumon M, Me´chin F, Pascault J-P. Polym Eng Sci 2002; 42(8):1724–39. [5] Meynie´ L, Fenouillot F, Pascault JP. Polymer 2004;45(6):1867–77. [6] Choplin L, Marchal Ph. Proceeding of XIIth international congress on rheology (1996), Quebec, Canada 1996 p. 665–6. [7] Aı¨t-Kadi A, Marchal Ph, Choplin L, Chrissemant AS, Bousmina M. Can J Chem Eng 2002;80:1166–74. [8] Cassagnau P, Fenouillot F. Polymer 2004: Doi 10.1016/j.polymer 2004.09.027. [9] Gimenez J, Cassagnau P, Michel A. J Rheol 2000;44(3):527–47. [10] Bouilloux A, Macosko CW, Kotnour T. Ind Eng Chem Res 1991;30: 2431–6. [11] Cassagnau P, Nietsch T, Michel A. Int Polym Process 1999;14(2): 144–51. [12] Gimenez J. PhD Thesis. University Lyon1; 1999. [13] Ferry JD. Viscoelastic properties of polymers, New York, 3rd ed Wiley & Sons 1980. [14] Bonnet A, Pascault JP, Sautereau H, Camberlin Y. Macromolecules 1999;32(25):8524–30. [15] Vinh-Tung C, Lachenal G, Chabert B, Pascault J-P. Toughened Plast 1996;2:59–74. [16] Venderbosch, R.W. PhD Thesis. Eindhoven University of Technology; 1995.
Rheological study of mixing in molten polymers: 2-mixing of reactive systems Philippe Cassagnaua,*, Franc¸oise Fenouillotb a
Laboratoire des Mate´riaux Plastiques et Biomate´riaux, UMR-CNRS 5627: Inge´nierie des Mate´riaux Polyme`res ISTIL, Universite´ Claude Bernard, Lyon 1, 15 Boulevard Latarjet, 69622 Villeurbanne cedex, France b Laboratoire des Mate´riaux Macromole´culaires, UMR-CNRS 5627: Inge´nierie des Mate´riaux Polyme`res, Institut National des Sciences Applique´es, 17 Avenue Jean Capelle, 69621 Villeurbanne cedex, France Received 12 March 2004; received in revised form 23 August 2004; accepted 15 September 2004 Available online 1 October 2004
Abstract Three reactive polymer systems have been examined with a new mixing device adapted on a classical rheometer in order to investigate reactive mixing situations encountered in polymer blends. After having characterized the bulk polymerization of 3-caprolactone (3-CL), the polymerization of 40 wt% of 3-caprolactone into a copolymer of ethylene and vinyl acetate (EVA) was run into the rheo-mixer. The kinetics of the reaction in dispersed media was observed slightly different from that in bulk since the characteristic time of 3-caprolactone diffusion into EVA is much lower than its time of mixing. On the other hand, it was observed that the molecular weight distribution of the poly(3caprolactone) is broader in dispersed media (Mw Z 15; 000 g moleK1 ; IpZ2.6) than in bulk (Mw Z 17; 000 g moleK1 ; IpZ1.6). A broadening of the molecular weight distribution in dispersed media was pointed out due to the fact that 3-CL monomer is partitioned between the EVA and PCL phases leading to a non-homogeneous concentration of monomer in the reactive phase. The polycondensation of 40 wt% of a epoxy-amine system into a polystyrene matrix was also investigated and the morphology of the resulting material examined. A gradient of structure and conversion was detected in a blend obtained from the assembly of two initially nonreactive layers. The gradient reveals that the amine diffuses faster than the epoxy leading to non-stoichiometry of the reactive functions across the sample. When the blend was polymerized under shear, the kinetic of the reaction remained unchanged regardless the level of shear. However, the morphologies were significantly different, pointing out the importance of the coalescence and droplet deformation phenomena. Spherical droplets were observed at 0.15 sK1, elongated droplets and fibers at 1.5 and 15 sK1. q 2004 Elsevier Ltd. All rights reserved. Keywords: Rheo-mixing; Diffusion; Reaction
1. Introduction The dispersion and/or dissolution at a molecular scale of low viscosity reactants in a polymer of high viscosity is a key for the control of the final macromolecular architecture when one considers a reactive polymer processing technology in the molten state. Such a reactive process in polymer blending can be for example the in situ polymerization of one phase in the presence of one other [1–3] or the grafting of a functional specie onto a polymer chain [4]. Generally
* Corresponding author. Tel.: C33 4 72 44 62 08; fax: C33 4 72 43 12 49 E-mail address: [email protected] (P. Cassagnau). 0032-3861/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2004.09.028
speaking, the time of mixing must be lower than the time of chemical modification and/or polymerization otherwise the local stoichiometry may be different from what is desired, resulting in non-controlled materials. For example, Cassagnau et al. [1] investigated from the torque variation in an internal mixer, the formation of polyethylene/thermoplastic-polyurethane blend (PE/TPU) via in situ polymerization of diols and diisocyanate in a molten polyethylene matrix. They demonstrated that the viscosity ratio, which considerably increases with the extent of the polymerization reaction, appears to be the physical parameter which governs the blend morphology development. In a same way, Meynie´ et al. [5] investigated the polymerization of a thermoset system into a polystyrene matrix. They show that
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the morphologies are strongly dependent, not only on the evolution of the viscosity ratio between the matrix and the dispersed phase, but also on the gel point and crosslinking phenomena of the dispersed phase. Nevertheless, these experiments were conducted in non-well controlled flow as the variation of the torque can be qualitatively correlated only to an apparent viscosity. On the other hand, due to the amount of polymer introduced in such mixer (about 60 g) and the low conductivity of polymer components, self heating phenomenon by viscous dissipation is another non-well controlled parameter. Reaction exothermicity may be another one. Depending on the processing conditions, i.e. apparent shear rate, a overheating of several degree is generally observed that accelerates the reaction. Finally, when mixing a large amount of a low viscosity reactant with a polymer in a internal mixer, the phenomenon of lubrication of the rotor by the liquid is often observed. In this case, the measured torque is apparently zero and no information on the state of mixing is available until the liquid slowly mix and/or react and stops lubricating the rotor. The rheoreactor concept provides an alternate solution to such rheological characterization problems [6]. The rheoreactor is a curved bottom vessel supplied with a helicoı¨dal double ribbon impeller. Nevertheless, such a system adapted to a conventional rheometer cannot be developed for polymeric system of high viscosity due to the torque upper limit authorized by the transducer of the rheometer. Consequently, this mixer-type rheometer can be only devoted to the study of intermediate viscosity liquids such as water/glycerol, salad dressing, polymer solutions or oligomers [7]. To overcome these difficulties, a mixing device called rheomixer has been developed [8] and adapted on a classical rheometer, with the aim of measuring reliable viscosity data on high viscosity fluids during the course of reactive mixing operations. In a previous work we have examined the process of mixing a low viscosity liquid, miscible in a high viscosity polymer [8]. The objective of the present study is to investigate, the influence of the mixing conditions on the polymerization kinetics and the morphology development of molten-reactive dispersed media. For this purpose, we analyze two systems: the polymerization of 3-caprolactone in a copolymer of ethylene and vinyl acetate (EVA) matrix and a the polycondensation of a thermoset (epoxy-amine) in a polystyrene (PS) matrix. For both, we expect to collect new data with the aim of better understanding the phenomena of diffusion, mixing and morphology formation.
2. Experimental 2.1. Experimental set-up The rheo-mixer has been described elsewhere in more details [8]. The rheo-mixer device was designed in order to
carry out the mixing experiments on a Rheometrics Mechanical Spectrometer (RMS800). This mixing device is based on a parallel plate analogy with some vertical obstacles like spikes in the gap for breaking up the streamlines of the fluid. A calibration procedure was used to convert the rotational speed/torque in shear/stress for a quantitative analysis of the torque curves. Most of the tests were conducted in the step shear mode. The viscosity was calculated from the torque values and the apparent applied shear rate according to the calibration constants [8]. Consequently, the mixing curves show the variation of the viscosity as a function of the time of mixing. Some tests were run in dynamic mode in the domain of the linear viscoelasticity. For the latter, a parallel plate geometry with 25 mm diameter was used. 2.2. Materials 2.2.1. Polymers A copolymer of ethylene and vinyl acetate (EVA) with a melt index of 420 was supplied by Atofina. The amount of acetate groups contained in this copolymers is 28 wt%. This copolymer was named EVA28420. Its zero shear viscosity is 360 Pa s at TZ100 8C. The polystyrene (PS) Lacqrene 1450 N was also supplied by Atofina. 2.2.2. Monomers and initiator 3-caprolactone (3-CL) and titanium tetrapropoxide were obtained from Aldrich and Roth Sochiel, respectively. The epoxy was a diglycidyl ether of bisphenol A (DGEBA) with a degree of polyaddition of nZ0.15 supplied by Bakelite. The diamine was 4,4 0 -methylenebis[2,6-diethylaniline] (MDEA) supplied by Lonza. 2.3. Sample preparation Globally, we applied the same experimental procedure described in a previous paper [8]. The particular procedures used for the reactive systems are detailed in the following section. 2.3.1. 3-Caprolactone The initiator was added to the 3-caprolactone just before the experiments in the rheometer. The reactive mixture was transferred, through a 2 ml syringe, to the vessel of the rheomixer preheated to the experimental temperature. The exchange surface between the reactive mixture and the area was very small in comparison with the volume, limiting the loss of monomer by evaporation. 2.3.2. EVA/3-caprolactone The EVA disc was brought into contact with the lower plate of the vessel at room temperature. The oven was then heated to the experimental temperature, TZ102 8C. The reactive liquid (3-CLCinitiator), was transferred in the vessel, through a syringe, over the EVA molten layer. The
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Fig. 1. Bulk polymerization of 3-caprolactone in the rheo-mixer at different shear rates. Rheo-kinetics at TZ102 8C.
polymer and liquid thicknesses were calculated in order to adjust the concentration of poly(3-caprolactone) in EVA at about 40 wt%. The total thickness of the reactive medium (EVAC3-CL) was 1.6 mm. 2.3.3. PS/thermoset The composition of the blend studied was 60 wt% of PS and 40 wt% of epoxy/amine thermoset precursors. In order to avoid a reaction of the monomers during a blending operation, two non-reactive pre-blends were prepared by extrusion. Their composition were calculated to obtain a stoichiometric mixture of epoxy functions with the amine functions of the co-monomer. The thickness of the nonreactive blends discs was adjusted in order to respect the epoxy-amine stoichiometry and the global composition of the blend. Thus: hPS/DGEBAZ0.9 mm and hPS/MDEAZ 0.7 mm. The epoxy rich disk was placed on the lower plate (25 mm parallel plate geometry) with the amine rich disk just above it. The total thickness of the reactive medium was 1.6 mm.
3. Results and discussion 3.1. Bulk polymerization of 3-caprolactone Modeling the change in viscoelastic properties during a reaction of polymerization is of great importance for predicting the flow behavior since flow is strongly dependent on the degree of polymerization. In a previous study [9], the rheological behavior during the bulk polymerization of 3-caprolactone with titanium tetrapropoxide as the initiator was studied directly in the rheometer with parallel plate geometry and in the domain of the linear viscoelasticity. Thus, the low deformation applied could not promote mechanical mixing. The aim of the present work is to study the influence of high deformations and mixing
conditions on the rheo-kinetics of 3-CL bulk polymerization. The length of the polymer chains depends directly on the ratio [M]/[Io]. [M] and [Io] are, respectively, the monomer and initiator concentrations. The polymerization can be classified as living because of the linear variation of the average molecular weights in number and in weight versus the conversion. For the present work [M]/[Io]z600. Fig. 1 shows that a same evolution of the viscosity (rheokinetics) during the polymerization is obtained for shear rates in the range 1.5–150 sK1. The viscosity changes from the viscosity of the 3-CL monomer at the early stage of polymerization to the viscosity of the PCL at the end of the polymerization. However, the description of the viscosity is less accurate for low shear rate, especially at the early stage of polymerization due to the sensitivity of the rheometer sensor. Besides, it can be pointed out that the PCL samples completely polymerized exhibit a Newtonian behavior over the whole range of shear rate investigated here. Furthermore, the molecular weights of the poly (3-caprolactone) samples at the end of the polymerization process in the rheomixer were measured by size exclusion chromatography system (SEC). For the different processing conditions, we observed the same molecular weights of the samples within an experimental error of 10%: Mw Z 17; 000 g moleK1 and Mn Z 10; 500 g moleK1 that is a polymolecularity index of IpZ1.6. Consequently, the fact that the viscosity–kinetic curves depict a same curve at different shear rate is not surprising since the system is thoroughly mixed before its introduction into the mixer. In other words, the polymerization process is carried out in a perfect homogeneous media. Such a result was already reported in bulk polymerization of linear polyurethane in extruder [10,11]. 3.2. Polymerization in dispersed media New developments in reactive polymer blends have
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Fig. 2. Polymerization of 3-caprolactone in the rheo-mixer during mixing with EVA28420 at different shear rates. TZ102 8C. Initial conditions: EVAC3CLZ1.6 mm. EVA/PCL at the end of the mixing and polymerization processes: 60/40 wt%.
increased sharply because reactive compounding may provide viable mechanisms for the in situ elaboration of the desired blend with controlled structure and morphology. However, the complex nature of polymer reactive modification processes in dispersed media requires a fundamental understanding of the mechanisms that govern chemical reaction in polymeric melt phases, as well as a critical assessment of the role of key process parameters on the endusage properties of the blend. The aim of this part is to investigate experimentally the effect of mixing on reaction kinetics and morphology development of reactive polymeric blends produced by polymerization of monomers in dispersed media and in controlled processing conditions (shear rate, temperature). For this purpose, two systems were considered. The first one is the formation of a EVA/PCL blend via in situ polymerization of 3-CL and the second one is a thermoplastic (TP)/thermoset (TS) blends resulting from the in situ polycondensation and of thermoset precursors in a polystyrene phase. 3.2.1. EVA/3-caprolactone We used the same experimental conditions as those used in bulk polymerization (TZ102 8C, [M]/[Io]z600) for the polymerization of the 3-CL in the molten EVA28420 matrix. Fig. 2 shows the variation of the apparent viscosity versus the mixing time (or reaction time) for different shear rates. Remind that the characteristic time of the bulk polymerization of 3-CL is 200 s as observed in Fig. 1. In dispersed media, Fig. 2 shows that the time of polymerization, tp, of 3-CL in EVA48420, is around 300–400 s, that is close to the time observed in bulk. In a previous work [8], we observed for the same processing conditions (temperature and shear rate) that the characteristic time, tm, for mixing 3-CL monomer with EVA28420 was around 3000 s. The comparison of these characteristic times indicates that
the polymerization rate is not really slowed down when the 3-CL polymerizes in presence of the EVA. Thus, the monomer does not completely incorporate to EVA prior to reacting. This is not very surprising since 3-CL and EVA are not miscible in all proportions as was observed visually and may be predicted from their respective solubility parameters (dCLZ21.1 MPa1/2 and dEVAZ18 MPa1/2). Furthermore we demonstrated, that the mixing process is controlled by the diffusion process of 3-caprolactone monomer in molten EVA phase [8]. By analogy with the Deborah number, we can define a dimensionless number as the ratio of the characteristic time of the mixing (tm) to the characteristic time of the polymerization process (tp): Db Z
tm tp
(1)
Here Db z10 confirming that the 3-CL polymerization is mainly performed in a bulk state. Indeed, Fig. 2 shows that the time of mixing/polymerization is in the same order of magnitude whatever the shear rates. The difference in the shape of the curves is due to the fact that the viscosity of EVA28420 and the viscosity of the PCL at the end of the polymerization are in the same order of magnitude (h0,EVA28420Z280 Pa s and h0,P3-Cl Z360 Pa s at TZ _ 1:5 sK1 Þ only a slight 100 8C). At low shear rate ðgZ variation in the viscosity curve is observed. This curve depicts an inflexion point around 100 s. At high shear rate _ 15 and 150 sK1) a slight lubricant effect is observed at (gZ 20 s followed by an increase of the viscosity due to the 3-CL polymerization. Actually, the viscosity ratio considerably changes with the extent of the polymerization, varying from l!10K5 at the early stage of the polymerization and mixing processes to lz1 at the end of these two coupled processes. The
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Fig. 3. Molecular weight distributions of poly(3-caprolactone) synthesized in bulk and in EVA dispersed media (1.5 and 15 sK1).
mixing between EVA and PCL is accelerated due this evolution of the viscosity ratio with the extent of the reaction of polymerization. This discussion is based on the fact that these two polymers are immiscible which has been proved by differential scanning calorimetry study of the present system at the end of the polymerization process. The controlling effect of the viscosity ratio on mixing was noted on other systems as the reactive blending by in situ polymerization of an polyurethane phase in a polyethylene media for instance [1,11]. Fig. 3 shows that the molecular weight distribution of the PCL is broader in dispersed media (Mw Z 15; 000 g moleK1 ;
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IpZ2.6) than in bulk (Mw Z 17; 000 g moleK1 ; IpZ1.6) and that the molecular weight distribution is not dependent on the mixing conditions (shear rate) as observed in Fig. 3. Actually, the PCL and EVA phases are PCL and EVA rich phases containing 3-caprolactone since the monomer is present until the end of the reaction, and can diffuse in EVA. Thus, the broadening of the molecular weight distribution in dispersed media can be explained by the fact that 3-CL monomer is partitioned between the EVA and PCL phases leading to a concentration gradient of the 3-CL in the reactive phase. We should also point out that the initiator of the 3-CL polymerization is also a catalyst of the transesterification reaction so that exchange reactions could be encountered between vinyl acetate groups and 3-CL ester groups. And actually, when the polymerization of 3caprolactone in EVA was carried out in extruder at 160 8C, NMR analysis have shown that 3-caprolactone was grafted onto EVA chains by a transesterification reaction [12]. Nevertheless, the polymerization temperature of the present work is low (TZ1008C) compared to the temperature in extruder (TO160 8C), and the reaction of transesterification can be neglected as was confirmed by NMR analysis. 3.2.2. PS/thermoset This part concerns the polymerization of a epoxy-amine thermoset system (DGEBA–MDEA) into a PS matrix. The
Fig. 4. PS/thermoset (60/40) reactive blending at 177 8C and 60 rpm in a internal mixer. (a) Glass transition of the PS phase (&) glass transition of the epoxy phase (:) and torque (—) as a function of time. Final morphology (b) reaction in an oven and (c) reaction in the mixer.
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Fig. 5. PS/thermoset (60/40). Modulus of the complex shear viscosity at uZ1 rad sK1 and TZ177 8C. Linear viscoelasticity—curve A: initial condition, two layers constituted by PS/DGEBA (hZ0.9 mm) and PS/MDEA (hZ0.7 mm); curve B: initial conditions, one layer constituted by a pre-blend PS/DGEBA– MDEA prepared in a batch mixer at TZ80 8C.
concentration of the thermoset precursors is 40 wt%. This system is quite complex and more details can be found in the work of Meynie´ et al. [5]. In the cited work, the authors considered the evolution of the blend under shear in a batch mixer at TZ177 8C. This study was based on a homogeneous initial blend of epoxy and amine monomers in PS, at least the authors supposed that the concentration in reactants at a micro-scale is at the stoichiometry. Their results confirmed this assumption. At the temperature of the experiments (TZ177 8C), the monomers are soluble into the PS so that the epoxy and amine functions are diluted and the reaction is slowed down compared to a pure DGEBA– MDEA mixture at the same temperature. When the chain length of the growing epoxy-amine species increases, phase separation occurs. A dispersed phase is formed by nucleation and growth. This phase is mainly composed of epoxy-amine oligomers and monomers. This causes an acceleration of the polymerization. Then, the composition of the phases evolves as can be seen from the Tg data, and the epoxy droplets grow in size by diffusion of the monomers and coalescence until the epoxy domains gel and the conversion goes to completion (Fig. 4(a)). When the reaction is run in quiescent conditions, the final particles are spherical (Fig. 4(b)) while under flow, in a internal mixer for instance, they deform and coalesce ending to large, irregular particles (Fig. 4(c)). Note that at the end of the process both the PS-rich phase and the epoxy phase are purified since their Tgs are those of the pure PS and pure DGEBA–MDEA network. In the present work, we aimed to study the influence of the diffusion and of the mixing of epoxy and amine monomers on the morphology development of the reactive blend with PS. A first experiment was run in dynamic regime at uZ 1 rad sK1. The strain amplitude was adjusted from 100% at
the beginning of the test to 1% at the end of the experiment to obtain a measurable torque and to remain within the domain of linear viscoelasticity. In linear viscoelasticity the entities (polymer chains, droplets,.) are deformed around their equilibrium position. In other words, the material structure is not altered during the course of measurement in the domain the linear viscoelasticity [13]. Thus we assume that this test was performed in ‘static’ condition, that is no coalescence was produced by the shear. This assumption is valid since we found that the morphology obtained with the operating conditions described above was similar to the morphology obtained in a oven in the absence of any shear (compare Fig. 6 with Fig. 4(b)). This demonstrates that we can assimilate experiments run in the rheometer in oscillatory mode, to experiments run in truly static conditions. Two different preparations of the sample test were experienced. The first one is a two layers sample constituted of a PS/DGEBA layer and a PS/MDEA layer. Each layer is not reactive unless it is mixed or put in contact with the other one. The second one is a unique layer of a homogeneous PS/DGEBA–MDEA blend prepared at low temperature (80 8C) in order to limit the reaction. We can see the evolution of the complex viscosity of the two different preparations of the samples in Fig. 5 showing the variation of the complex shear viscosity versus reaction time at TZ177 8C. Curves A and B in Fig. 5 exhibit an identical qualitative variation of the viscosity with polymerization time. Note that the viscosity data brings information on the evolution of the polymerization and on the morphology of the blend [14–16]. The phase separation can be well identified for both curves in Fig. 5. When the epoxy phase separates, the PS matrix viscosity rises. It is the first change of the slope of the viscosity variation with time. Later, the gelation is identified by a marked increase of the viscosity.
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Fig. 6. Condition for: 1-two initial layers PS/DGEBA and PS/MDEA (curve A, Fig. 5); (a) sample from PS/DGEBA layer external side, (b) sample from interface between PS/DGEBA and PS/MDEA layers and (c) sample from PS/MDEA layer external side, and 2-one initial layer constituted by one homogeneous PS/DGEBA-MDEA layer (curve B, Fig. 5).
We find the phase separation around 22 min, and the gel point is detected around 37 min. Finally, the rheology of the blend above the gel point of the thermoset phase is essentially governed by the rheology of the PS rich matrix since the thermoset phase has a droplet morphology. The increase in viscosity above the gel results from the migration of amine and epoxy monomers and/or oligomers out of the PS-rich phase toward the epoxy domains. Consequently, the viscosity and the glass transition temperature of the PS matrix increases with the extent of the reaction. The final viscosity in curve A is 10 times lower than the viscosity in reference curve B indicating that the epoxyamine reaction is globally slower, and that the epoxy conversion is not 100% at the end of the process. For the two layers sample, it is clear that the diffusion of the monomers
in the neighbor layer is controlling the process. At the beginning, we note that the polymerization is only slightly slower since the phase separation and the gel are detected, respectively, at 22 and 42 min. This indicates that in this period of time, the diffusion is faster than the polymerization. Latter, the evolution of the viscosity, thus of the epoxy conversion, is much slower. The decrease of the diffusion coefficient of the monomers may be the cause since the global viscosity and glass transition temperature of the phases increase with the conversion of the epoxy, making the polymerization diffusion controlled. The final morphology of PS/DGEBA–MDEA from curve B (one layer) consists of a dispersion of spherical thermoset rich particles of about 3 mm diameter (see Fig. 6.2). A similar structure has been observed by Meynie´ et al. [5] when the polymerization occurred in static conditions.
_ 0:15 sK1 and TZ177 8C. Curve A: initial condition, two layers constituted by PS/DGEBA (hZ Fig. 7. PS/thermoset (60/40) blending in rheo-mixer. gZ 0.9 mm) and PS/MDEA (hZ0.7 mm); curve B: initial conditions, one layer (hZ1.6 mm) constituted by a pre-blend PS/DGEBA–MDEA prepared in a batch mixer at TZ80 8C.
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Fig. 8. Scanning electron microscopy of PS/thermoset blend at the end of _ 0:15 sK1 Þ at TZ177 8C. (a) from the curing test in dynamic condition ðgZ two initial layers PS/DGEBA and PS/MDEA. (a) From initial homogeneous PS/DGEBA–MDEA system: curve B in Fig. 7. (b) From two initial layers PS/DGEBA and PS/MDEA: curve A in Fig. 7.
Morphology analysis in Fig. 6 shows that a different result is observed for the droplet morphology of the initial two layers PS/DGEBA–PS/MDEA. For the analysis of these morphologies, it must be kept in mind that the epoxy and amine reactants are initially separated so that the reaction may happen only if the monomers diffuse to the adjacent layer. We can assume that diffusion occurs essentially in one direction (axial direction of the sample). The morphology at the external side of the initial PS/DGEBA layer consist in hard thermoset droplet with an average diameter about 5 mm (Fig. 6.1(a)). In the region, initially at the interface between the PS/MDEA and PS/DGEBA layers (Fig. 6.1(b)), droplets with smaller size about 2 mm are observed. The morphology at the external the side of the initial PS/MDEA layers (Fig. 6.1(c)) consists in sub-micron particles with an average diameter about 0.2 mm. This structure is typical of a late nucleation showing that the epoxy concentration was low in
this zone of the sample so that the epoxy droplets did not have time enough to grow in size. From the above remarks, we can deduce that the diffusion of the amine in the PS/ epoxy layer is much faster than the diffusion of the epoxy in the PS/amine layer. Consequently, the DGEBA and MDEA reactants are not at the stoichiometric balance over the sample thickness. Moreover, it can be assumed from this gradient of morphology, a gradient in phase compositions leading to particles with different glass transition temperature. On the other hand, the non-stoichiometric condition implies that some DGEBA and MDEA monomers and/or oligomers remains in the PS phase (in the PS/amine layer) at the end of the polymerization process. As a consequence the conversion is not complete and this explains why the viscosity in curve A is 10 times lower than the viscosity in reference curve B after 80 min of reaction. Fig. 7 shows the variation of the absolute complex viscosity versus polymerization/mixing time at a low shear _ 0:15 sK1 Þ: Curves A and B were obtained from the rate ðgZ same initial conditions as early described in Fig. 5. For an initial homogeneous reactive blend, Fig. 8(a) shows a homogeneous morphology as expected, with droplets diameter of about 10 mm. The droplets remain spherical under weak dynamic conditions but their diameter is larger than that under static conditions: 3 mm (Fig. 6.2). The balance between the mechanism of coalescence and break up is shifted to the coalescence process and leads to larger particles. If we come back to Fig. 7, before the gel, the same evolution of the viscosity is observed compared with Fig. 5. However, after the gel, the viscosity of the two layers sample (curve A) reaches the same value than in curve B at the end of the polymerization (hz105 Pa s). According to our previous discussion, this result means that above a certain viscosity, even a very low shear rate is efficient enough to induce a mechanical mixing that decreases the distance of diffusion of the epoxy to such an extend that the
Fig. 9. PS/thermoset blending in rheo-mixer at different shear rate, TZ177 8C. Initial conditions: one layer (hZ1.6 mm) constituted by a pre-blend PS/DGEBA–MDEA prepared in a batch mixer at TZ80 8C.
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Fig. 10. PS/thermoset blending in rheo-mixer at different shear rate, TZ177 8C. Initial conditions: two layers constituted by PS/DGEBA (hZ0.9 mm) and PS/MDEA (hZ0.7 mm).
polymerization kinetics becomes almost identical to that of the homogeneous sample. Nevertheless, the viscosity is a macro-scale measure as the whole volume of the sample is solicited in the experiment. Morphological analysis in Fig. 8(b) shows that the morphology, resulting from mixing of initial two layers, is not homogeneous at a micro-scale level. From this figure two distinct zones can be observed, the first one is constituted by thermoset droplets with 10 mm diameter and the second one is constituted by smaller droplets with a size lower than 1 mm. The size of these zones is approximately 100 mm. From these results, we conclude that the rheo-mixer working at a low shear rate is efficient to get a macro-scale mixing (scale about 100 mm) for the present system. Figs. 9 and 10 show the variation at high shear rates _ 1:5 and 15 sK1) of the absolute complex viscosity (gZ versus the time of mixing/polymerization for initial homogeneous PS/DGEBA–MDEA blend and initial two PS/DGEBA and PS/MDEA layers, respectively. Just after the phase separation process the viscosity decreases
Fig. 11. Morphology of PS/thermoset blend at the end of the curing test in dynamic condition from initial homogeneous PS/(DGEBA–MDEA) system. For morphology at 0.15 sK1 see Fig. 8(a); (a) 1.5 sK1 and (b) 15 sK1.
depending on the shear intensity. At the gelation process _ 15 sK1 it rapidly the viscosity suddenly increases. For gZ drops few minutes later. Actually the sample becomes _ 0:15 sK1 the unstuck from the mixer plate. Above gZ microscopy analysis (Figs. 11 and 12) shows a rough morphology with a fibrous structure. Some thermoset ellipsoids of 50 mm length can be observed in these figures. The coalescence and deformation of the droplets has produced an elongated structure. Finally, viscosity curves and morphology pictures do not show any noticeable difference between both initial conditions (homogeneous reactive system and non-homogeneous two layers) so that the mixing efficiency of the rheo-mixer is rather good for two initial liquids with a viscosity ratio close to one.
4. Conclusion The reactive mixing process of three different systems has been investigated with the help of a new mixing device. Reactive systems such as bulk 3-caprolactone polymerization and dispersed media polymerization of linear thermoplastic (poly 3-caprolactone), and thermoset (epoxy/amine) were investigated. We showed that the polymerization in EVA molten polymer is slightly slowed down at high shear rates by a dilution effect resulting from the mixing of 3caprolactone in EVA phase. It was observed that the molecular weight distribution of the poly(3-caprolactone) is broader in dispersed media (Mw Z 15; 000 g moleK1 ; IpZ 2.6) than in bulk (Mw Z 17; 000 g moleK1 ; IpZ1.6). This broadening of the molecular weight distribution in dispersed media was pointed out due to the fact that 3-CL monomer is partitioned between the EVA and PCL phases leading to a non-homogeneous concentration of monomer in the reactive phase.
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the Scanning Electron Microscopy work and Sophie Dumas, PhD student, for her help in the polymerization of the 3-caprolactone. The authors also acknowledge one of the reviewers for their very useful comments.
References
Fig. 12. Morphology of PS/thermoset blend at the end of the curing test in dynamic condition from two layers constituted by PS/DGEBA (hZ 0.9 mm) and PS/MDEA (hZ0.7 mm). For morphology at 0.15 sK1 see Fig. 8(b); (a) 1.5 sK1 and (b) 15 sK1.
For the PS/thermoset system, we have used two initially non-reactive layers of PS/DGEBA and PS/MDEA. The experiment run in oscillatory mode, that is in the absence of mechanical mixing allowed to show that the apparent polymerization rate was progressively slowed down in the two layer system. This was attributed to decrease of the diffusion coefficient of the monomers caused by the variation of the glass transition temperature of the phases. Besides, the amine was found to diffuse faster than the epoxy leading to a gradient of morphology and conversion in the two layer sample. The morphology of the blends polymerized under increasing levels of shear depicted spherical thermoset droplets only at very low shear rate. Above 0.15 sK1, epoxy elongated droplets and fibers were produced.
Acknowledgements The authors gratefully acknowledge Pierre Alcouffe for
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