B Vieille
email: benoit.vieille@insa-rouen.fr

L Taleb

Keywords: A, Fabrics -A, PMCs -B, High-temperature properties -C, Notch

Could thermoplastic-based composites be used to replace thermosetting-based composites in high-temperature secondary aircraft structures ? The purpose of this work is to establish the ability of a material system to be used in aircraft engine nacelles when subjected to static loadings, with a key upper temperature of 120°C. In order to provide answers to this question, the thermo-mechanical behaviors of carbon fiber fabric reinforced PPS or epoxy laminates have been compared specifically within the temperature change with 120°C at the upper bound. The temperature-dependent ductile behavior of laminates is more or less exacerbated, depending on polymers' glass transition temperature, and laminates' stacking sequence. For both materials, the degree of retention of tensile mechanical properties is quite high in notched and unnotched quasi-isotropic laminates. A Digital Image Correlation technique has been used in order to understand the influence of temperature and matrix ductility on the mechanisms of overstresses accommodation near the hole. In fabric reinforced laminates, the hightemperature results suggest a competition between the mechanisms of damage, and the mechanism of plasticization, enhanced in angle-ply lay-ups. Thus, the highly ductile behavior of TP-based laminates, at temperatures higher than their T g , is very effective to accommodate the overstresses near the hole.

INTRODUCTION

Thermosetting (TS) resin-based composite materials have been extensively used over the past 40 years. Even though they display a high degree of mechanical properties retention at high temperature, they also present undeniable drawbacks, such as the need for low-temperature storage, a hard-to-control reticulation process, a very long curing process, and handmade draping, which causes most of the irreversible defects of the manufacturing process. High-performance thermoplastic (TP) resins, such as polyetheretherketone (PEEK) and polyphenylenesulfide (PPS), offer a promising alternative to TS resins such as epoxies: high degree of chemical resistance, excellent damage and impact resistance, and they may be used over a wide range of temperatures.

Unlike TS resins, TP resins may be melted down after they are formed, thus providing recycling possibilities and different manufacturing. However, TP composites are more difficult to repair than conventional TS composites, and high-performance TP are more expensive. PPS resin is very thermally stable, it has fire retardant properties, it can attain a high degree of crystallinity (almost 60%) and it is cheaper than PEEK resin [START_REF] Myers | Stress-state effects on the viscoelastic response of PPS-based thermoplastic composites. Advances in TP Matrix Comp. Mat[END_REF].

The basic idea of this work is to find out whether it is possible to replace thermosettingbased composites used in nacelles of plane's engines or not. The service temperature of these nacelles, that usually consist of C/Epoxy laminates, is 120°C. The question was therefore to evaluate how detrimental such a degree of temperature can be on the behavior of PPS-based laminates compared to epoxy-based laminates. The physical properties (Crystallinity rate and glass transition temperature T g ) of TS and TP resins are closely associated with the mechanical behavior of PMCs under high temperatures [START_REF] Spruiell | A review of the measurement and development of crystallinity and its relation to properties in neat PPS and its fiber reinforced composites[END_REF]. An overview of environment effects on performances of PMCs (TS and TP) is proposed in [START_REF] Hancox | Overview of effects of temperature and environment on performance of polymer matrix composite properties[END_REF]. Around their T g , the nonlinear behavior of fiber-reinforced polymer composites becomes significant, especially under off-axis loading conditions. This nonlinear response, associated with the shear deformation of the polymer matrix along reinforcing fibers, is enhanced at high temperature depending on the matrix viscoplastic -3 nature [START_REF] Kawai | High-temperature off axis fatigue behaviour of unidirectional CFR composites with different resin matrices[END_REF][START_REF] Lafarie-Frenot | Comparative in-plane shear behaviour of longcarbon-fibre composites with TS or TP matrix[END_REF][START_REF] Deng | The non-linear response of quasi-isotropic composite laminates[END_REF][START_REF] Ha | Modeling the viscoplastic behavior of fiberreinforced TP matrix composites at elevated temperatures[END_REF][START_REF] Jen | Mechanical properties in notched AS4/PEEK APC2 composite laminates at elevated temperature[END_REF][START_REF] De Baere | Modeling the nonlinear shear stressstrain behavior of a CFR PPS from rail shear and [45,-45] 4S tensile test[END_REF]. Very few studies have investigated the high-temperature behavior of PPSbased laminates [START_REF] Meyer | Temperature effect on reinforced TP composite properties for primary aircraft structure applications[END_REF][START_REF] Vieille | Influence of temperature on the behavior of carbon fiber fabrics reinforced PPS laminates[END_REF][START_REF] Loverich | Life of PPS composites subjected to cyclic loading at elevated temperatures. Time dependent and nonlinear effects[END_REF][START_REF] Mahieux | Elevated temperature bending stress rupture behavior AS4/APC2 and comparison with AS4/PPS[END_REF][START_REF] Miyano | Loading rate and temperature dependence of flexural fatigue behavior of a satin woven CFRP laminate[END_REF] or woven-ply epoxy laminates [START_REF] Abot | In-plane mechanical, thermal and viscoelastic properties of satin fabric carbon/epoxy composite[END_REF][START_REF] Kawai | Off-axis fatigue behavior of plain weave carbon/epoxy fabric laminates at room and high temperatures and its mechanical modeling[END_REF][START_REF] Osada | Initial fracture behavior of satin woven fabric composites[END_REF].

The fact of joining composites to other structures using mechanical fasteners (bolts or rivets) makes it necessary to drill laminates, which damages both fibers and matrix in the area around the hole. But notches have a significant impact on the growth of damaged zones and on the mechanical properties of laminates. The notch sensitivity, defined by the ratio of laminates' notched and unnotched ultimate strengths, depends on its ability to accommodate overstresses near the hole. This ability, associated with matrix ductility, implies specific damage mechanisms for TP-and TS-based laminates.

A good review about the hole sensitivity of TS-based UD laminates is proposed in [START_REF] Awerbuch | Notched strength of composite laminates: Predictions and experiments -A review[END_REF], and a few studies have reported the tensile behavior of notched C/TP (PEEK or PPS) laminates [START_REF] Carlson | Notch sensitivity of thermoset and thermoplastic laminates loaded in tension[END_REF][START_REF] Lagace | Response of notched C/epoxy and C/PEEK systems[END_REF][START_REF] Touchard-Lagattu | Damage and inelastic deformation mechanisms in notched TS and TP laminates[END_REF][START_REF] Touchard | Mechanical behaviour characteristics of a TP composite used in structural components[END_REF][START_REF] Aymerich | Response of notched AS4/PEEK and C/epoxy laminates subjected to tension fatigue loading. Fatigue and Fracture of Eng[END_REF]. C/PEEK TP composite laminates are more sensitive to the presence of holes than TS laminates. A large amounts of damage have been observed in notched TS composite specimens because of a brittle matrix. Such extensive damages have not been observed in C/PEEK notched laminates because of a much more ductile matrix. In lay-ups with a large percentage of 0° plies, the presence of larger damage zones at notches in C/epoxy, compared to C/PEEK, permits a reduction of the stress concentration leading to a higher failure load [START_REF] Brillaud | Experimental determination by laser speckle patterns of plastic strains near a hole in a TP composite laminate[END_REF]. Most of the attention has been focused on UD ply laminates even though fabric laminates tend to be more resistant to splitting and delamination than similar UD laminates [START_REF] De Azevedo Soriano | Notch sensitivity of C/epoxy fabric laminates[END_REF]. The occurrence of splitting in UD laminates is facilitated by the low strength in the transverse direction, while delamination can occur due to both Poisson's ratio and shear coupling mismatch. In balanced fabric laminates, 45° plies are self-balanced and delamination due to shear coupling mismatch is limited. However, both woven-ply and UD-ply C/epoxy exhibit similar notch sensitivity [START_REF] Lagace | Notch sensitivity of graphite/epoxy fabric laminates[END_REF]. The effects of matrix ductility (TS and TP) on unnotched/notched carbon woven-ply laminates has been studied in [START_REF] Kawai | Effects of matrix ductility and progressive damage on fatigue strengths of unnotched / notched CF plain woven roving fabric laminates[END_REF]: the failure -4 mechanisms of C/epoxy laminates are exclusively governed by the progressive delamination from notch tips, while no appreciable effects are observed on the notch tip fracture of C/TP laminates. To the authors' knowledge, most of the studies on the behavior of notched/unnotched polymer-based composites (TS or TP) have dealt with tests at room temperature. There are no significant results about the high temperature behavior of high-performance TP-based laminates. There is no reference about the plastic accommodation of overstresses, specifically observed in angle-ply notched laminates subjected to tensile loadings.

MATERIAL AND METHODS

Materials

The composite materials studied in this work are carbon fabric reinforced PPS [START_REF] Vieille | Influence of temperature on the behavior of carbon fiber fabrics reinforced PPS laminates[END_REF] or epoxy prepreg laminate plates. The toughened PPS resin (Fortron 0214) and the epoxy resin (914) are respectively supplied by the Ticona and by Hexcel. The woven-ply prepreg , supplied by SOFICAR, consists of 5-harness satin weave carbon fiber fabrics (T300 3K 5HS). The volume fraction of fibers is 50% in C/PPS and C/epoxy laminates.

The prepreg plates are hot pressed according to two lay-ups: [0/45/0/45/0/45/0] and [45] 7 . Even though the number of effective plies is higher in 0°/90° plies than in 45° plies , the fabric being balanced in warp and weft directions, the first lay-up can almost be considered Quasi-Isotropic (Q-I), and is suitable for applications in aircraft engine nacelles. In Angle-Ply (A-P) laminates [45] 7 , each ply consists of fibers oriented at +45° (warp direction), and at -45° (weft direction). Even though an A-P lay-up is not suitable for use in aircraft structures, it allows the authors to investigate specifically the plastic accommodation mechanism in order to understand the specific contribution of TP matrix to the high-temperature behavior of high-performance TP-based laminates.

Method

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All the quasi-static tests were performed using a 100 kN capacity load cell of an MTS 810 servo-hydraulic testing machine at room moisture. The temperature control system (oven and temperature controller), provided a stable temperature environment during the test. Tensile tests were conducted on notched and unnotched laminates with two stacking sequences. Their mechanical properties were determined according to the EN 6035 (Tensile -Open hole tensile tests) and EN 6031 (In-plane shear test) standards.

Strains are evaluated from an extensometer and V denotes constant crosshead speed (in mm/min) applied to the specimen during displacement-controlled tests. Four specimens were tested in each configuration (See Fig. 1). The mechanical properties were averaged from these values (See Tab. 3 and Tab. 5). The average thickness of specimens is about 2.22mm in C/PPS, and about 2.08mm in C/epoxy.

In accordance with the previous standards, the longitudinal and shearing modulus were calculated from the following definitions :

Where is the difference in the applied tensile loads at and .

S is the specimen cross section.

represents the shearing distortion.

(Difference in the longitudinal strains) and (Difference in the transversal strain) are calculated from the values given by strain gages located along X and Y axis (See Fig. 1).

RESULTS

A Differential Scanning Calorimetry performed by the CRISMAT laboratory on C/PPS gave T g =98°C and a cristallinity rate equal to 30%, and T g =190°C for C/Epoxy. Tensile tests have been performed in the direction along the 0° axis. The modulus and strength in warp and weft directions are assumed to be equal (See Tab. 1). For balanced woven fabrics, it is reasonable to admit such an equality for different composite systems [START_REF] Naik | Prediction of on-axes elastic properties of plain weave fabric composites[END_REF].

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However, such an assumption also depends on the geometry of textile fabric (satin or plain weave). Indeed, it is only verified when the crimp ratio, which represents the undulation of weft fiber bundles over warp fiber bundles, is low e.g. for satin weave fabric [START_REF] Osada | Initial fracture behavior of satin woven fabric composites[END_REF]. In this latter case, the equality of the elastic properties has been observed experimentally in balanced 5-harness satin weave carbon fiber fabrics reinforced epoxy laminates [15][29]. In a work dealing with the laminate mechanics for balanced woven fabrics [START_REF] Akkerman | Laminate mechanics for balanced woven fabrics[END_REF], Akkerman showed that this assumption is also verified in balanced 5harness satin weave carbon fiber fabrics reinforced PPS laminates.

Tensile tests on unnotched laminates

In Q-I laminates, according to the CLT (Classical Laminate Theory) calculations, the load carried by 0° plies (57% of the plies) represents 73% of the total load. Thus, the laminate response is mainly dominated by the behavior of 0° fibers resulting in a temperature-independent behavior in C/Epoxy and C/PPS laminates (See Fig. 2a and2b). The dotted line corresponds to the slope at the origin of the curve. It represents a purely elastic response, and emphasizes a behavior characterized by a very low degree of non-linearity. Load/unload tensile tests performed on C/Epoxy laminates (See Fig. 3a) show that, even at 95% of the load at failure, the longitudinal modulus of the laminate is still very high in comparison with its initial value at both temperatures (See Tab. 2a). It suggests that this slightly non-linear behavior is not associated with the gradual failure of 0° oriented fibers during loading, but with a localized plastic behavior of 45° oriented plies. However, such a behavior is limited by the 0° plies and plasticity cannot develop. At 120°C, the ductile behavior of the plain PPS resin is enhanced, but the laminate response being mainly fiber-dominated, then the laminate behavior is not significantly influenced by the matrix response at high temperature (See Fig. 4a). For design requirements, it appears that the tensile properties of Q-I C/epoxy laminates are greater than C/PPS ones (see Tab. 3). This is amplified as temperature increases around -7

the T g of C/PPS composites. In A-P laminates, the matrix-dominated response results in a elastic-ductile behavior (See Fig. 5a and5b) in both composite systems. From a general standpoint, the in-plane shear properties (modulus and strength) are higher in C/Epoxy than in C/PPS laminates (Tab. 3). Load-unload tensile tests performed on A-P C/Epoxy laminates show a gradual decrease of material's longitudinal modulus (See Fig. 4b ) as well as residual strain (See Tab. 2b). These observations highlight two simultaneous mechanisms (damage and plasticization) that will be discussed in the next section. Contrary to the Q-I lay-up, the highly ductile behavior of PPS can be observed in A-P lay-ups (See Fig. 4a), more particularly at 120°C because the test temperature is higher than its T g . Thus, the shear modulus decreases more in C/PPS (-67% ) than in C/Epoxy laminates (-21%). Both materials have virtually the same strength at both temperatures, but the elongation at failure is a lot more significant in C/PPS laminates (+420% at RT and +367% at 120°C). The CLT and the Tsai-Hill failure criterion gives the failure strengths with an acceptable relative change (See Tab. 4).

The in-plane shear stress-strain response of both materials consists of different regimes.

The initial one is elastic. It is followed by a non-linear region which began with the onset of matrix plastic deformation, and which was fully developed at shear strains of 3-6% in C/epoxy and 17-27% in C/PPS. The stress borne by the composite increased linearly until failure afterwards, as the displacement compatibility between matrix and fibers led to the progressive rotation of the latter with the applied shear strain and increased the axial stresses borne by the fibers [START_REF] Totry | Effect of fiber, matrix and interface properties on the in-plane shear deformation of CFR composites[END_REF]. In epoxy-based laminates, the low ductile behavior of epoxy matrix does not allow the fibers to rotate a lot (About 5° at

RT and 7° at 120°C -See Fig. 6). The development of interface fracture significantly modified the in-plane shear behavior as it hindered the load transfer between matrix and fibers and precipitated fracture by the coalescence of interface cracks. The observation of the fracture surface indicates that the specimen failure is associated with the onset of -8

a shearing band oriented at 45° (See Fig. 6) in reference to the loading direction [START_REF] Karayaka | Deformation and failure behavior of woven composite laminates[END_REF]. In the case of PPS-based laminates, the highly ductile behavior of PPS matrix limits the development of interface fracture, and promotes the rotation (About 17° at RT and 26° at 120°C -See Fig. 6). Such mechanisms cause necking that is clearly observable on the C/PPS stress-strain curves, as well as on the failed specimens. Necking comes along with an important delamination. The magnitude of the maximum rotation of fibers bundles can be evaluated in unnotched laminates at both temperatures (See Fig. 6). Such an assessment of fibers rotation rests on the observation of the fibers at the specimen surface. It is therefore not very accurate in comparison to the evaluation of the rotation and the progressive development of damage done by optical microscopy [START_REF] Totry | Effect of fiber, matrix and interface properties on the in-plane shear deformation of CFR composites[END_REF].

According to the conclusions done in [START_REF] Kawai | Off-axis fatigue behavior of plain weave carbon/epoxy fabric laminates at room and high temperatures and its mechanical modeling[END_REF][17] for woven-ply composites, and depending on matrix nature, cracks appear within fiber bundles and in the area where the weft fibers undulate over warp fibers. It results in the splitting of the initially perpendicular weft and warp fibers bundles.

3.2Tensile tests on notched laminates

The hole factor represents the hole sensitivity of the laminate ultimate strength. It is defined by where and are respectively the notched and unnotched ultimate stresses, t and w being the thickness and width of the specimens. From these definitions, the tensile responses can be compared for both materials at both temperatures (Fig. 2 and Fig. 5). In Q-I laminates, temperature influences very little both strength and stiffness of both materials (See Tab. 5). The hole factor increases very little bit showing that hole sensitivity is almost temperature-independent in fiberdominated laminates. Thus, a service temperature equal to 120°C is not significantly detrimental for the tensile behavior of notched laminates with a Q-I lay-up. In A-P -9

laminates, the ductile behavior of the polymer matrix is more or less enhanced at high temperature depending on the matrix. It contributes to adversely affect the mechanical properties of both materials. However, they become less hole-sensitive as temperature increases (See Tab. 5). C/PPS is less resistant (about -20%) than C/Epoxy, and elongation at failure in C/PPS is much more pronounced than in C/Epoxy.

DISCUSSION

Knowledge

Digital Image Correlation technique for strain field measurements

Full-field measurement techniques using Digital Image Correlation are an effective method in investigating the behavior of heterogeneous materials such as composite structures. This technique is of great interest for the detection of damage zones, and for quantifying their increase, particularly, in high gradients zones. Strains fields around holes can be measured by DIC or optical systems using laser speckle pattern method [24][34-35]. The ability of DIC to measure full-field in-plane surface deformations at high temperatures was evaluated in [START_REF] Lyons | High-temperature deformation measurements using digital-image correlation[END_REF]. The DIC technique consists in painting the sample white, and then in sputtering black paint on the surface of the sample in order to get a painting with different grey value patterns. When a displacement is applied to the sample, a Region Of Interest around the hole is filmed with a CCD video camera, whose -10

resolution is 659*493 pixels, and which is equipped with a zoom lens. DIC is based on the comparison of specimen's images acquired at 10Hz by the software Visionstage.

Two images are chosen at two specific strain states : before the test starts and just before failure. The correlation software VIC2D tracks blocks of pixels in order to measure surface displacement and build up full field 2D strain maps. The hole being a singularity at the surface of the specimen, the blocks of pixels located on the notch boundary cannot be used for calculating the strain field. Thus, the offset in the hole's radius represents the area where the calculation is not possible. The strain maps built gives the in-plane Green-Lagrange strain field. The local accuracy of these values depends on many parameters : fineness and density of the sputtered black painting on the white surface, the grid step, the pattern size and the heterogeneity of laminates. The use of a DIC technique at high temperature requires to film the specimen through a porthole' glass, making necessary to avoid catch lights on the porthole's glass that might be detrimental for the accuracy of calculated strain fields. For each lay-up, the comparison of strain components E YY (Y being the loading axis) at both temperatures reflects the influence of matrix ductility on the mechanisms of overstresses accommodation near the hole via changes in overstrains fields. The overstrain coefficient can be defined as the ratio of the maximum strain near the hole over the strain applied to the specimen.

Quasi-isotropic laminates

In woven-ply C/Epoxy laminates, notch-tip blunting (splitting) is known to reduce the stress concentration effect of the notch, and then to limit the laminate's hole sensitivity in Q-I laminates [19][20]. Contrary to UD-ply laminates [START_REF] Cowley | Damage accumulation at notches and the fracture stress of carbon-fibre/polymer composites: Combined effects of stress and temperature[END_REF][34], there is no clear sign of delamination in woven-ply C/Epoxy laminates [START_REF] Kawai | Effects of matrix ductility and progressive damage on fatigue strengths of unnotched / notched CF plain woven roving fabric laminates[END_REF], where 45° plies are selfbalanced, and delamination due to shear coupling mismatch is limited. In such materials, transverse and longitudinal cracks appear early during tensile tests. They extend fast and widely around the hole, whereas in C/TP laminates, damage and plastic -11

strains remain localized in the hole vicinity [8][10]. According to X-ray observations [START_REF] Lagattu | High strain gradient measurements by using digital image correlation technique[END_REF], splitting appears later and is less extended than in C/Epoxy, resulting in a higher notch sensitivity of C/TP [27][33-35]. In the present work, Figure 3b shows that the residual strains after unloading are very small (less than 0.07%) in notched laminates for both materials. It suggests that plasticity cannot develop in the 45° plies of in Q-I laminates, even at 120°C when the polymer matrix ductile behavior becomes more significant. When the 0° fibers break, the load is transferred transversely to the adjacent fibers that progressively fail until final failure of laminates. Thus, the fracture surfaces is characterized by a transverse failure mode (See Fig. 11).

In C/Epoxy laminates, fracture surfaces are similar across the range of test temperatures.

At RT, the profile of the fracture surface is typical of brittle fracture with almost no bare fibers and no hackles (See Fig. 11a), and almost no delamination as the edge views of failed notched specimens can suggest it (See Fig. 12a). At 120°C, grooves can be clearly observed in the fracture area between plies what means fracture is dominated by matrix cleavage (See Fig. 11b). The tensile responses and fracture surfaces of Q-I C/Epoxy laminates suggest that damage mechanisms do not seem to change significantly as temperature increases. It results in very similar longitudinal strain fields at both temperatures (See Fig. 7). Observation of stress-strain curves suggests a catastrophic failure of notched specimens as in the unnotched case (See Fig. 2).

Observations of fracture surface in C/PPS laminates show bare fibers in hackles configuration (See Fig. 11a), confirming the relatively weak fiber/matrix bond that was observed in unnotched laminates [START_REF] Mahieux | Elevated temperature bending stress rupture behavior AS4/APC2 and comparison with AS4/PPS[END_REF]. At 120°C, as the number of bare fibers has increased, it is reasonable to conclude that interfacial bond has degraded and interfacial failure has become extensive (See Fig. 11b), resulting in some delamination visible on the edge views of failed notched specimens (See Fig. 12a). Thus, the longitudinal strain fields are different at RT and 120°C (See Fig. 8).

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In Q-I laminates, the strain concentration increases significantly with increasing temperature for both materials (See Tab. 6), but in the meantime, the hole factor slightly increases (See Tab. 5). It suggests that the effect of notch is locally enhanced by the increase in temperature whereas the notch sensitivity (being the reverse of the hole factor) slightly decreases for both materials. On the one hand, the ductile behavior of both matrix systems is enhanced at 120°C, but the matrix behavior is limited by the Q-I lay-up. On the other hand, the strength of the fiber/matrix bond and the interlaminar shear strength decrease as temperature increases. It results in a little more extended damage near the hole (splitting) which is combined with a more extended plastic behavior at high temperature. It explains why the strain concentration increases with temperature. From the overstresses accommodation standpoint, these mechanisms contribute to a little release of stresses in the bundles of fibers near the damaged /plasticized area. Thus, the hole sensitivity slightly decreases, but the C/PPS mechanism associated to the ductile nature of PPS is slightly less effective than C/Epoxy one.

Angle-ply laminates

The tendency is different from the one observed in the Q-I laminates because the mechanical behavior of A-P laminates is mainly matrix-dominated. Consequently, they are less hole-sensitive, more particularly at 120°C (See Tab. 3 and 5). The ability of the matrix to transmit load to fibers becomes impaired at high temperature, but plasticity is more or less exacerbated depending on the matrix nature. While matrix undergoes a plastic deformation along the 45° oriented fibers, these fibers tend to rotate (See Section 3.1) . Because of the undulation of weft fibers over warp fiber bundles, weft fibers are overstressed and fail first, resulting in failure surfaces oriented along the warp direction. Even though matrix ductility has an impact on stresses distribution around the hole through a plasticity mechanism, it is not the only mechanism for -13 accommodating overstresses. In C/PPS laminates, cracks along 45° fibers and some fiber/matrix debonding can be observed in the 45° direction to the hole as well as delamination (See Fig. 12b). The longitudinal strain fields are very different at both temperatures (See Fig. 10), because plasticity is prominent in PPS-based laminates along the 45° oriented fibers. It finally results in a lower strain concentration near the hole as temperature increases (See Tab. 6). It also suggests that there is competition between the mechanism of damage and the mechanism of plasticization enhanced by a temperature increase. Such mechanisms contribute to release the stresses near the hole, hence a better notch sensitivity at high temperature. In C/Epoxy laminates, the tensile responses (See Fig. 2b) show that the ductile behavior is slightly enhanced at 120°C what contributes to moderate changes in damage mechanisms as temperature increases.

The longitudinal strain fields are very similar at both temperatures (See Fig. 9), and the overstrain coefficients are virtually the same (See Tab. 6). The observations of failure surfaces show interfacial debonding in warp and weft directions (stairs profile) along the 45° direction, and some delamination (See Fig. 12b). These mechanisms combined with a larger plastic deformation along the 45° oriented fibers cause the increase of C/Epoxy laminates' hole factor with increasing temperature. Still considering the detrimental effect of a temperature increase on the strength of the matrix, of the fiber/matrix bond and interlaminar shear strength, damage appears earlier than at RT.

Both mechanisms (plasticization and damage) counter act, resulting in virtually the same strain concentration. Thus, the hole sensitivity is almost temperature-independent

Figure captions

  of the extent of strains at stress concentrators like holes is important in engineering design because it is associated with the magnitude of the localized stresses at the discontinuity, and hence the notched strength of the structure. In order to study the influence of temperature and matrix ductility on the mechanisms of overstresses accommodation near the hole, a Digital Image Correlation technique has been used. It gives qualitative and quantitative information about strains fields. The following discussion is based on the damage mechanisms in woven-ply notched laminates reported in[START_REF] Touchard-Lagattu | Damage and inelastic deformation mechanisms in notched TS and TP laminates[END_REF][START_REF] De Azevedo Soriano | Notch sensitivity of C/epoxy fabric laminates[END_REF][27] at RT, as well as under high-temperature[START_REF] Jen | Mechanical properties in notched AS4/PEEK APC2 composite laminates at elevated temperature[END_REF][16][START_REF] Akkerman | Laminate mechanics for balanced woven fabrics[END_REF].

  in A-P epoxy-based laminates. Lastly, these mechanisms are very effective because the ultimate notched strength is equal to 82% of the unnotched laminate's strength. The prominent plastic accommodation observed at high temperature in notched C/PPS laminates remains less effective than C/Epoxy one, due to less degree of stress relieving damage formation around the notch.Changes in the mechanical behavior of notched/unnotched laminates were investigated in carbon fabric reinforced PPS and epoxy laminates between RT and 120°C.Considering the results of the present study, PPS-based composites can be used to replace epoxy-based composites in high-temperature secondary aircraft structures subjected to static loadings. Despite the maximum test temperature (120°C) is higher than its T g (98°C), the degree of retention of mechanical properties at elevated temperature is very high in both notched and unnotched C/PPS laminates with a Q-I layup. Far from epoxy's T g , a high temperature has also very little influence on both notched and unnotched tensile properties of C/Epoxy laminates. In A-P laminates, changes in mechanical properties and failure mechanisms are significant in both materials even though the epoxy matrix ductility at 120°C is less significant than the PPS matrix one. From the overstrain coefficient, it is possible to evaluate the efficiency of the plastic accommodation mechanism to limit overstresses near the hole. It is salient that a temperature increase has a positive impact on the hole sensitivity of both materials (+34% in C/PPS and +9% in C/Epoxy laminates). The issue of TS vs TP composites was first addressed 20 years ago, and since then, it is known that the ability of a material to accommodate overstresses near the hole at RT is less effective in C/TP than in C/Epoxy laminates, due to less degree of stress relieving damage formation around the notch. At high temperature, the contribution to knowledge base from this work rests on the emphasis of plastic accommodation mechanisms near the hole in both warp and weft directions. Such mechanisms are specific to woven-ply TP-based laminates at temperatures higher than their Tg, when the TP matrix behavior becomes highly ductile. These mechanisms are effective to release overstresses resulting in lower hole-sensitivity. At last, it is still necessary to investigate the long-term behavior of both notched/unnotched woven-ply laminates over the same range of temperature, before -15 such materials can actually be considered for applications in aircraft engine nacelles. A further work will deal with these issues in order to understand how the specific accommodation mechanisms near the hole and the corresponding damage accumulation change over time due to environmental factors and cyclic (or creep) loading.

Fig. 1 -Fig. 2 -Fig. 3 -Fig. 4 -Fig. 5 -Fig. 6 -Fig. 7 -Fig. 8 -Fig. 9 -Fig. 10 -Fig. 11 -Fig. 12 -

 123456789101112 Fig. 1 -Geometry of tensile specimens: (a) notched -(b) unnotched Fig. 2 -Tensile vs open hole tensile tests at both temperatures on Q-I laminates: (a) C/PPS -(b) C/Epoxy Fig. 3 -Load/unload tensile tests on Q-I laminates: (a) unnotched C/Epoxy -(b) notched C/Epoxy and C/PPS at 120°C Fig. 4 -Comparison of mechanical responses: (a) Q-I and A-P C/PPS laminates vs plain PPS resin -(b) Load/unload tensile tests on A-P C/Epoxy laminates Fig. 5 -Tensile vs open hole tensile tests at both temperatures on A-P laminates: (a) C/PPS -(b) C/epoxy Fig. 6 -Rotation of fibers bundles for both materials: (a) RT -(b) 120°C Fig. 7 -Quasi-isotropic C/Epoxy laminates at RT and 120°C. Comparison of failed specimens (transverse failure in weft direction) and corresponding Green-Lagrange longitudinal (warp direction =y) strain fields obtained by Digital Image Correlation: the longitudinal overstrain near the hole increases with temperature (+40%) Fig. 8 -Quasi-isotropic C/PPS laminates at RT and 120°C. Comparison of failed specimens (transverse failure in weft direction) and corresponding longitudinal (warp direction = y) strain fields obtained by Digital Image Correlation: the longitudinal overstrain near the hole increases a lot with temperature (+158%) Fig. 9 -Angle-ply C/Epoxy laminates at RT and 120°C. Comparison of failed specimens (shearing band along the warp direction) and corresponding longitudinal strains fields obtained by Digital Image Correlation: the longitudinal overstrain near the hole is virtually the same at both temperatures (+6%) Fig. 10 -Angle-ply C/PPS laminates at RT and 120°C. Comparison of failed specimens (failure along warp and weft directions at RT -failure along warp direction at 120°C) and corresponding longitudinal strains fields obtained by Digital Image Correlation: the longitudinal overstrain near the hole decreases a lot with temperature (-52%) due to extended plastic deformation along the 45° oriented fibers Fig. 11 -Quasi-isotropic notched failed specimens: (a) Side macro views of fracture surfaces highlight the transverse failure mode -(b) Alternative views showing the specific fracture mechanisms in each material Fig. 12 -Edge views of notched failed specimens: (a) quasi-isotropic -(b) angle-ply

Tab. 3 -Tab. 4 -

 34 Tensile mechanical properties (individual test values and average) at both temperatures for both materials: quasi-isotropic and angle-ply laminates Calculated vs experimental failure strengths: quasi-isotropic and angle-

  11.05-11.11-11.17-11.24: Changes in overstrain coefficients for both materials

						0.68-0.71-	2.4-2.42-
					11.14	0.72-0.73: 0.71	2.43-2.51: 2.44	-5	+356
	Influence of		(MPa)	-3	-8	+6	-18
	temperature	C h		+4.0	+34	+11	+9
	(%)		(%)	+3	+238	+3	+38
					C/PPS	C/Epoxy
					Q-I	Angle-ply	Q-I	Angle-ply
			(%)		0.64	3.29	0.63	1.72
	RT		Max. overstrain (%)		0.9	14.25	1.07	5.75
			Overstrain coefficient	1.41	4.33	1.70	3.34
			(%)		0.66	11.14	0.71	2.44
	120°C	Max. overstrain (%)		2.4	23.3	1.66	8.65
			Overstrain coefficient	3.64	2.09	2.37	3.54
	Influence of temperature on the overstrain coef. (%)	+158	-52	+40	+6
		Tab. 6 -

Tab. 5 -Open-hole tensile mechanical properties (individual test values and average) at both temperatures for both materials: quasi-isotropic and angle-ply laminates

ACKNOWLEDGEMENTS

This work is a part of the TOUPIE project, and is supported by the DGE through the MOV'EO competitiveness cluster, in which several partners are collaborating (Région Haute-Normandie, Aircelle, AMPA, AXS Ing., Université du Havre, ENSI Caen, ENSM Paris, INSA Rouen).


Miruna Dorobantu Bodoc

Laurent Prat

Catherine Xuereb

Christophe Gourdon

Thierry Lasuye

Keywords: Microreactor, Free Radical Polymerization, On-line Monitoring

On-line Monitoring of Vinyl Chloride Polymerization Reaction in Microreactor using Raman Spectroscopy

A novel capillary-based microfluidic device has been designed to follow the vinyl chloride polymerization reaction. The use of a co-flow generation system enabled obtaining monodisperse vinyl chloride droplets within 200 µm in diameter, each one being considered as a polymerization reactor. During polymerization VCM droplets were visualized with a high speed camera. At the end of the reaction PVC grains were observed with a Scanning Electron Microscopy technique. Real-time non-invasive Raman measurement has been performed on stationary vinyl chloride monomer droplets and has provided values of effective reaction orders n and effective rate constants k. This microdevice allowed the investigation in difficult conditions (pressure, temperature) with a minimal amount of reagents and consequently under safe conditions.

Introduction

Microreactor technologies present several advantages compared to their macro-scale counterparts. For instance, due to their high surface-to-volume ratio (up to 4000 times more important than the conventional batch reactors) the specific exchange surfaces and the mass and heat transfer rates are substantially greater. This implies an enhanced heat control permitting to perform highly endothermic or exothermic reactions. The confined and laminar flow allows precise control over the size of the dispersed phase or the residence time. Another obvious advantage of the miniaturization is the small volume of reagents involved enabling to test dangerous reactions otherwise difficult to perform at laboratory scale. Various applications of microreactors have increased in recent years [1,2]. Among the latter, new methods for accurately obtaining and controlling polymer particles with narrow size distribution have been described in the literature [3 -6]. Polymerization in a two-phase dispersion system, such as suspension reactions, is the main method for obtaining polymer particles and several related studies have been conducted on this subject. Various authors proved droplet milifluidics to be an appropriate device for polymerization, further adding analytical tools which allow reaction monitoring [7]. The suspension polymerization of vinyl chloride (S-PVC) by a free-radical mechanism consists in dispersing liquid monomer under its autogeneous pressure into small droplets in a continuous aqueous phase containing surfactants and initiator. The protective colloids added inhibit the droplets to coalesce. The reaction starts by the thermo-decomposition of the organic initiator into the vinyl chloride monomer (VCM). As a consequence, the reaction proceeds in monomer droplets, which may be considered each as a bulk reactor. Because PVC is insoluble in its own monomer, VCM polymerization is a heterogeneous process involving several physical transitions. The final PVC product is composed of primary particles and their agglomerates. The kinetic monitoring of suspension polymerization is a challenging task due to the high heterogeneity of the reaction medium.

Different techniques are present in the literature for the measurement of conversion during this type of reaction. Usually this parameter is obtained by gravimetry and only a relatively small number of works involve monitoring the suspension polymerization with techniques such as ultrasounds [8], calorimetry [9], NIR spectroscopy [10] or Raman spectroscopy [11]. Although considered as a common reaction, the S-PVC has never been studied at the micro scale. This might be due to the particularities of the operatory conditions (pressure, temperature) and to the toxic nature of the monomer. In this context the microtechnologies offer promising perspectives for the synthesis of PVC, since they imply generation of droplets with the same size, shape and chemical composition. The safety of the process would be granted by the small quantities of VCM involved and by the facility to evacuate the reaction heat. Raman spectroscopy seems as an appealing method for non-invasive kinetic investigation of suspension polymerization reaction. This technique offers information on both chemical structure as well as concentration within microdroplets, allowing real time tracking. Recent studies present the successful coupling to microfluidic devices in polymerization reactions [12]. However, literature does not reveal any research of this analysis method applied to S-PVC. The main objective of this study was to implement a device allowing on-line monitoring of the S-PVC reaction on an isolated droplet in microchannel. This was accomplished by coupling a non-invasive Raman system enabling quantitative analysis of monomer/polymer droplets in microreactor. Also real time particle visualization was performed by connecting a high speed camera to the microchannel. First vinyl chloride monomer droplet behaviour images were therefore obtained.

Materials

Vinyl chloride monomer (VCM) was supplied by Ineos ChlorVinyls, France in a 500 mL pressure cylinder (Swagelok). Partially hydrolyzed poly(vinyl acetate) (PVA) used as suspending agents were Gohensol KP08, 72.5% hydrolysis and Active 45, 45% hydrolysis and were provided by Nippon Gohsei Japan and respectively AKZO. Bis(tert-butylcyclohexyl) peroxydicarbonate (Perkadox 16W) used as initiator for the polymerization reaction was received from Sigma Aldrich. The experimental set-up used for the VCM polymerization is represented on Figure 1. The microreactor test bench was designed to allow a maximum working pressure of 30 bars and a 90°C working temperature. The components exposed to VCM, such as tubing or valves, obtained from Swagelok are all made of stainless steel. Water phase and VCM flows are ensured by neMesys high pressure syringe pumps equipped with high pressure stainless steel syringes. The reactor consists of fused silica capillary tube. The tube where the reaction is carried out was rolled up on a stainless steel support which allows guiding the capillaries in a precise order. This support was sunk in a water bath of controlled temperature by a thermostat. At the end of the microchannel a waste reactor was connected in order to recover any traces of unreacted VCM. Pressure is measured on VCM circuit and in the collecting vessel by pressure sensors with an accuracy of +/-0.15% (Serv Instrumentation). The water and VCM phase encounter in a co-flow generation device as shown in the zoom in Figure 1. The VCM generation is performed at the nozzle tip of a 50 µm ID (150 µm OD) fused silica capillary tube placed into another capillary tube of 250 µm ID (360 µm OD). Once droplets formed, they are flowing inside a 15 cm long microtube and they are ejected into a larger fused silica capillary of 520 µm ID (670 µm OD). A secondary water flow is injected. The system was chosen in order to ensure that droplets do not touch the microtube's walls, producing therefore flow instabilities. The capillary length varied from 140 -200 cm.

Sequence of operations

For the suspension polymerization, the entire circuit was first purged with nitrogen to avoid inhibition by oxygen. The suspending agents were added to the aqueous phase in quantities of respectively 90 ppm for Gohensol KP08 and 45 ppm Active 45 with respect to the water phase. The initiator was pre-dispersed into the VCM and then pushed into the microchannel in order to disperse the VCM phase into the water phase. The flow rates varied for the VCM phase between 0.2 -0.5 mL/h and for the water phase: 2 -5 mL/h. Once the hydrodynamic regime was established, a first Raman spectrum was acquired at ambient temperature serving as a reference measure. Afterwards the reactor is heated. As soon as the desired temperature is reached, the entire flow was stopped, the inlets and outlets are subsequently sealed and spectra is recorded at constant time intervals. The polymerization temperature is fixed at different values in a range between 58 to 70 °C. The pressure in the microreactor is between 10 to 16 bars. At the end of the experiment the circuit is cleaned with a PVC solvent and with the carrier fluid. Droplet visualization is performed with a binocular Nikon SMZ-10 coupled with a high speed camera HCC -1000 (VDS Vosskühler GmbH) acquiring up to 1800 images/second. Light is ensured by fiber optics placed underneath the microchannel. Images are transferred to PC and processed with NV 1000 software (New Vision Technologies). The spectra were obtained with a Raman spectrometer (RXN-1, Kaiser Optical System Inc., U.S.) characterized by a near-infrared laser diode (400 mW, 785 nm) as excitation source. The microfluidic device was placed under the probe head with non-contact sampling optics (10x) of the spectrophotometer. The data acquisition was carried out through a 5 mm window on the capillary channels. The probe head was focused above the microchannel by moving a y-z barrel micrometer. The acquisition time of each spectrum was of 120 seconds over 1 accumulation. This interval was chosen to give an acceptable signal-to-noise ratio for the acquired spectra. The experimental set-up was designed to allow droplet visualization during the reaction by means of a CCD camera. It is possible as well to recover from the microchannel the polymerized grains and analyse them with a Scanning Electron Microscope (SEM) technique. By way of example, a VCM drop was monitored with a CCD camera at the first instants of the polymerization. The images corresponding to this experiment are presented in Figure 2. A change of colour at the droplet interface was clearly observed. After 2 minutes the interface starts to become cloudy, at 5 minutes a certain change in colour was noticed from white to grey and at about 7 minutes the droplet is almost black. After 10 minutes, the entire interface became completely black. This change of colour was assigned to the formation of a skin around the droplet. The formation of the thin "skin" at the droplets interface is believed to take place at the very early stages of the polymerisation (about 2% conversion) due to the adsorption of the PVA on the VCM/water interface. The skin is believed to be formed of a graft copolymer (PVA-PVC) transforming the adsorbed stabiliser (PVA) into a rigid layer of about 10 nm thickness [START_REF] Allsopp | Manufacturing and Processing of PVC[END_REF][START_REF] Eliassaf | [END_REF]. The primary particles formed inside VCM droplets will aggregate at this 'skin' leading to the formation of a perricellular membrane of approximately 200 -500 nm in thickness [15,16]. The pericellular membrane is indeed present on the PVC particles obtained in microchannel at the end of the reaction and it was observed by means of SEM technique. The membrane's surface is presented in Figure 3a and its thickness was measured at 418 nm (Figure 3b). During polymerization in batch reactors, primary particles move to the inner surface of the droplets by centrifugal forces caused by the agitation. These primary particles adhere and as they grow further they will finally coagulate with increasing monomer conversion. Consequently, a rigid layer is formed at the inner surface of the droplet covered by a PVA-PVC crosslinked skin at the outer surface of the droplet. The thickness of this layer probably depends on the agitation speed. If a higher agitation speed is used the skin will become thicker since more primary particles will move towards the inner surface of the particles driven by the centrifugal forces. Following this approach, in quiescent conditions as those in microchannel the perricellular membrane should have been thinner. However, its dimensions are strongly dependent on the PVA's concentrations employed. In our experiments the PVAs concentrations were beyond the critical micelle concentration therefore one may suspect their influence on the membrane's thickness. The outside surface of the particles (PVA-PVC graft copolymer membrane) is relatively smooth as it may be noticed in both photos in Figure 3. The folds and hollows in the membrane are a consequence of the shrinkage of the monomer droplet during the reaction since the density of the PVC is greater than the one of VCM (1PVC = 1,4 g.cm -3 1MVC = 0,86 g.cm -3 ). The droplet's evolution in micro channel during polymerization was monitored for 5 hours at 60°C and respectively 70°C. The images in Figure 4 show the evolution of the same droplet at different moments in time. For clearer view on the volume change, the droplets' contour was outlined (white circle at instant 0 and at 5 hours). The measurements of diameters were carried by tracking the interface of the droplet in the capillary. The image processing software furnished the corresponding positions which were subsequently transformed from pixel units into micrometers by using the channel width as scale. For the volume calculations, droplets were considered as spheres. As expected, the droplet volume contracts due to the density difference between monomer and polymer. The volume decrease with time is presented in Figure 5a. We remark a decrease of about 30 % of the droplet volume during monitoring time.

Various authors [17], [18] evaluated some of the physical properties of the two-phase dispersion water/VCM. They took into consideration an expression for the volume contraction of the polymerizing droplet due to density changes. Both before mentioned authors suggest a droplet volume model that evolves with conversion X:

X V V ε - = 1 0 (1) p m p ρ ρ ρ ε - =
with V, V 0 the volumes respectively at conversion X and X=0, 1 the shrinkage factor, 2 m and 2 p the monomer density and polymer density respectively. This type of calculation allows predicting droplet conversion only by the means of the measures performed on the pictures recorded with the CCD camera. The results are presented in Figure 5b. Nevertheless, it is important to underline the degree of uncertainty associated to this method. Therefore the results are only used to estimate a polymerization characteristic time. PVC grains recovered from the microreactor are around 200 µm in diameter as observed in Figure 6a. SEM pictures in Figure 6b were taken through a cross-section of a PVC grain. Normally, when conversion increases, the volume of free VCM inside the droplets decreases. Thus, the primary particles inside the droplets increase in size and fuse together determining the porous character of the droplet. The size of the primary particle was measured at about 1 µm in diameter (Figure 6b, small circle), value in good agreement with the average size mentioned in the literature [20]. The strong capillary forces that appear within the droplet enable the primary particles to fuse together. The result is the formation of a continuous open network of agglomerates of about 3µm in diameter (Figure 6b, large circle).

First kinetic analysis

In this section the on-line kinetic monitoring of vinyl chloride in the microdevice is presented and discussed. Figure 7a shows the spectrum of monomer droplets taken as reference measure. Although the spectra were obtained on droplets in aqueous phase their quality is good enough and no fluorescence was detected. The Raman band specific of the vinylic group exhibits a strong signal at 1606 cm -1 . The low frequency at 395 cm -1 is attributed to the planar deformation mode of the C=C-Cl skeleton, the 707 cm -1 corresponds to the stretching of the C-Cl band, the responses in the 1200-1400 cm -1 area indicate the CH 2 planar deformation or rocking modes and the magnitudes in the 3000-3100 area are assigned to stretching modes of the C-H bonds. Once the heating is started the reaction is initiated in the monomer droplets. During polymerization, monomer content decreases with conversion. Therefore the corresponding vinylic C=C related peak area undergoes an important decrease according to time, as represented in Figure 7b. Important variations are observed especially within the acquisitions at the end of the reaction. At high conversions, when the monomer concentration diminishes, the area becomes rather small and therefore more affected by the signal noise. Figure 8a, illustrates the area under the vinyl stretching band as obtained from the spectra monitored over time. It can be noticed the decreasing tendency despite the severe signal distortions which might be due to the medium heterogeneity. Moreover, due to the heating, the capillary tube is subjected to dilatation, therefore moving out of focus. In both cases, the area of incidence of the Raman laser might be modified from a whole polymer particle, to a set of particles, a region containing both water and polymer or to a complete de-focusing. All these constraints make representative sampling more difficult. Therefore spectrum out of focus has been removed. Afterwards, a criterion based on experimental observations was applied in the absence of a calibration method. C-Cl bond stretching (present in both monomer and polymer) was used as an internal standard and the area ratio of the C=C stretching mode to the C-Cl stretch was used to create a calibration for monomer predictions. This method based on peak area ratios was applied to the experimental values presented in Figure 8a (raw data) therefore leading to the values from Figure 8b (treated data). It can be observed that the area ratio decreases relatively continuously with time. The monomer concentration was assumed to be directly proportional to the normalized area of the vinylic peak in Figure 8b. Several authors proposed complex kinetic models for the S-PVC including a two-phase polymerization scheme [19,20]. The objective of this study is to obtain basic kinetic data in order to validate the global approach of this work. Therefore a simple kinetic model has been proposed in order to represent the experimental data. Note that in our calculation we assumed an ideal mixing. Given the important volume variations of the polymerizing droplet with conversion, this parameter was taken into account in the kinetic expressions. The rate equation for irreversible reactions with variable volume was expressed as described in equation:

V kM dt VM d n = - ) ( (2)
where M is the monomer concentration, k is the effective rate constant and n the effective reaction order. Equation 1 was used for volume determination. This simple kinetic model was chosen to fit the experimental data and to characterize the results. Integrating the equation 2 between the limits one can have access to a linearized rate equation for n-order reactions:

[ ] [ ] ) 1 ( ) 2 ( 1 1 0 0 2 0 2 ε V M V kt n M M t n n t - - + = - - (3)
where [M] t is the monomer concentration at the moment t calculated as the area of vinylic peak normalized by the C-Cl peak, [M] 0 is the initial monomer concentration, V t corresponds to the polymerized droplet volume at the instant t and V 0 is the initial droplet volume. From the experimental data the coefficients n and k (equation 3) were determined using a fitting method applied in the Excel Solver. From the normalized peak areas, the molar conversion may be estimated with the following relationship:

) ( ) ( ) ( ) ( 1 ] [ ] [ 1 ) ( 0 0 0 t A t A t A t A M M t X CCl M CCl M - = - = (4)
where A(t) is the peak area at time t corresponding to monomer (M) or C-Cl bond (CCl) respectively.

The experimental conversion and the calculated one versus time are compared in Figure 9 for two temperatures, 58 and 66 °C respectively. The solid line stands for the kinetic model issued from the equation 4. For the two temperatures the values of the reaction order n and the effective rate constant k are summarized in Table 1. As expected, Figure 9 shows that the polymerization rate increases when the reaction temperature is increased. The simplified kinetic model proposed seems adapted to the experimental values. Several aspects need to be pointed out: at 66°C the model seems better adapted while for 58°C the evolutions of the two curves are not fully comparables. The experiments were ended when a threshold was attained. At that point we estimated the reaction was complete and therefore a conversion of 1 was assumed. Note the similar reaction times as those presented in Figure 5b. The kinetic results obtained with the volume contraction evaluation from The effective rate constants obtained allow an estimation of the apparent activation energy 3E for the S-PVC polymerization with the equation:

1 2 1 1 2 ln 1 1 k k R T T E - 1 1 2 3 4 4 5 6 - - = ∆ (5)
A value of 80 kJ/mol resulted from calculation. It seems in agreement with some values reported in literature: activation energy of propagation in styrene polymerization (39 kJ/mol) [21], or in synthesis of polyurethanes (150 kJ/mol) [22]. For the vinyl chloride suspension polymerization, T.De Roo [23] furnished data on the activation energy of 24.9 kJ/mol for propagation and 54.3 kJ/mol for chain transfer. The value obtained in microreactor lies in the same magnitude order.

Conclusion

The main objective of this study was to implement a device allowing the study of S-PVC kinetics by means of an on-line analysis inside an isolated droplet in microchannel. This was accomplished using the Raman spectroscopy as a technique for quantitative analysis of monomer/polymer droplets in microreactor. A qualitative approach was also put in place by means of visualization of the PVC particles during or at the end of the reaction. A microfluidic device was designed to allow high pressure and temperature conditions in a perfectly safe manner. The quantities of vinyl chloride monomer, which presents a highly toxic character, were minimised, therefore the process remained completely reliable at any time for the operator as well as for the ambient environment. A co-flow device was set up to enable monodisperse VCM droplet generation. An isolated monomer droplet was visualized during the reaction. Its volume contraction was evaluated in terms of reaction conversion. Fused silica capillaries allowed application of Raman analysis technique for monitoring of monomer droplets suspended in aqueous solution. Good quality data have been acquired using this spectroscopy method applied to small droplets of monomer. The relatively long acquisition time (2 minutes) is not a limiting factor for data acquiring because the monitoring is performed on a stationary droplet. After spectra normalization, the monomer conversion degree was estimated with an acceptable accuracy versus time. A first model of the polymerization kinetics was performed with satisfactory results. Further improvement of a mathematical calibration technique may increase the measurement accuracy. Also the development of a complex kinetic model is in progress. The latter is suited for confined S-PVC conditions as those in microreactor. At the end of the reaction PVC grains were collected away from the device and analysed. Their dimensions and shapes were in good agreement with published data from a morphological point of view.

Symbols

Figure 1 .

 1 Figure 1. Experimental set-up used for the vinyl chloride polymerization at microscale.

Figure 2 .

 2 Figure 2. A VCM droplet during the first 10 minutes of reaction

Figure 1 .

 1 Figure 1. SEM images of the surface of the membrane surrounding the PVC particles obtained in microchannel (a) SEM measure of the membrane's thickness (b).

Figure 4 .

 4 Figure 4. Droplet evolution during 5 hours of reaction at 60°C

Figure 5 .

 5 Figure 5. Polymerized droplet volume measurement (a), Monomer conversion estimated with equation 1 (b).

Figure 6 .

 6 Figure 6. PVC grains recovered from the microreactor at the end of the reaction (a), cross-section of a PVC grain from the microreactor at the end of the reaction (b).

Figure 7 .

 7 Figure 7. Vinyl chloride monomer spectra at the initial time of reaction (a), and vinylic group area evolution with time (b).

Figure 2 .

 2 Figure 2. Raw vinylic group stretching area evolution during 250 minutes at a temperature of 58°C (a), Vinylic group stretching area normalized to the C-Cl stretch band area during 250 minutes at 58 °C (b).

Figure 9 .

 9 Figure 9. Kinetic curves obtained at 58°C and 66°C.

Table 1 . Reaction orders and rate constants at different working temperatures.

 1

Acknowledgements

The authors acknowledge Ineos ChlorVinyls for having funded this study.