Mechanical properties of woven glass fabric reinforced in situ polymerized poly(butylene terephthalate) composites

COMPOSITES SCIENCE AND TECHNOLOGY Composites Science and Technology 67 (2007) 390–398 www.elsevier.com/locate/compscitech

Mechanical properties of woven glass fabric reinforced in situ polymerized poly(butylene terephthalate) composites Z.A. Mohd Ishak

a,*

, Y.W. Leong b, M. Steeg c, J. Karger-Kocsis

c

a

c

School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia, 14300 Nibong Tebal, Penang, Malaysia b Advanced Fibro Science, Kyoto Institute of Technology, Sakyo-ku Matsugasaki, 606-8585 Kyoto, Japan Institut fu¨r Verbundwerkstoffe GmbH, Technische Universita¨t Kaiserslautern, Erwin-Schroedinger Str, Geb 58, D-67663 Kaiserslautern, Germany Received 19 December 2005; received in revised form 17 July 2006; accepted 12 September 2006 Available online 1 November 2006

Abstract Cyclic butylene terephthalate (CBT) has been polymerized in situ at T = 190 C in the presence and absence of woven glass fabric (WGF). The thermal properties of the in situ polymerized poly(butylene terephthalate) (ISP-PBT) were compared with those of a commercial injection molded PBT (IM-PBT). It was found that the crystallinity of IM-PBT is markedly lower than that of ISP-PBT rendering the latter more brittle. WGF-reinforced (ca. 50 vol.%) ISP-PBT composites were fabricated by compression molding at T = 190 C using displacement and pressure-controls. Effects of the molding condition were investigated in tensile, three-point bending, short beam shear and dynamic mechanical thermal analysis tests. Both tensile and flexural properties (stiffness and strength), as well as inter-laminar shear strength, were enhanced when the molding occurred under pressure-controlled instead of displacement-controlled conditions. Scanning electron microscopic (SEM) inspection revealed excellent wet-out of the fibers and good interfacial bonding between the fibers and ISP-PBT.  2006 Published by Elsevier Ltd. Keywords: Cyclic oligomers; Cyclic butylene terephthalate; Continuous fiber polymer composites; Thermoplastic composites; Poly(butylene terephthalate)

1. Introduction Considerable efforts have been devoted on combining thermoplastic resins such as polypropylene (PP), polyamide (PA), polyetherimide (PEI), polyphenylene sulphide (PPS), polyethersulphone (PES) and polyetheretherketone (PEEK) with various forms of continuous fiber reinforcements (e.g., unidirectional, woven fabrics) to produce thermoplastic composite materials [1,2]. One of the obvious targets was to replace thermoset based composites which have monopolized various applications viz. aerospace, automotive and sports industries. Thermoplastics offer a number of important advantages over thermosetting resins. *

Corresponding author. Fax: +6 4 594 1011. E-mail addresses: zarifi[email protected] (Z.A. Mohd Ishak), [email protected] (Y.W. Leong), [email protected] (M. Steeg), [email protected] (J. Karger-Kocsis). 0266-3538/$ - see front matter  2006 Published by Elsevier Ltd. doi:10.1016/j.compscitech.2006.09.012

Low scrap/good recyclability, better toughness/damage tolerance, unlimited shelf-life, and rapid fabrication cycle are some of their attractive features. On the contrary, thermoset composites suffer several drawbacks such as brittleness, lack of post forming and welding, poor recyclability, and often even moderate process efficiency. However, one of the main issues that are frequently raised in using thermoplastics as polymer matrices is their melt viscosities which are in the range of 100–10,000 Pa s. This creates serious problems in wetting and impregnation of reinforcing fibers, resin flow during fabrication and removal of entrapped air. One approach, which seems to be a promising route, is to develop thermoplastic resins which have the ability to be polymerized reactively like thermosetting resins but possesses the properties of thermoplastics. This has been achieved recently via the introduction of cyclic oligomers. Cyclic oligomers, such as cyclic butylene terephthalate

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(CBT), have a low processing viscosity which can be as low as 0.02 Pa s (water-like). Since cyclic oligomers may be transformed into linear high molecular weight thermoplastics via entropically-driven ring-opening polymerization in a short time scale (
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able PBT, namely Ultradur B4520, were supplied by BASF (Ludwigshafen, Germany). This material, designated as IM-PBT, served as reference in order to compare the properties of IM-PBT and in situ polymerized PBT (ISP-PBT). 2.2. Composite processing ISP-PBT and woven glass fabric reinforced PBT (WGFPBT) samples were prepared by in situ polymerization of CBT compacted by a compression molding press, SATIM Constructeur (Rion des Landes, France), equipped with temperature and pressure-controlled plates. Laminates containing 14 plies of WGF were molded at T = 190 C by using a stainless steel mold cavity of 150 · 190 · 3 mm3. CBT powder was dispersed between the WGF layers. To avoid any risk of moisture the prepared fiber package was kept in an oven at 90 C for 1 h prior to be transferred in the mold. The mold was placed at the center of the press the programmed cycle of which has been automatically started. Two approaches have been adopted in preparing the WGF-PBT composites. One involved displacement control whereby the specimens were molded by closing the two heating-plates under a prescribed condition as highlighted schematically in Fig. 1a. The composite manufactured under displacement-controlled (DC) condition is designated as WGF-PBT (DC). The mold was heated to 190 C. Care was taken to ensure that pressure was not applied to the non molten rigid powder in order to prevent fiber undulation. A stepping downwards (Xt) from about 10–3 mm was set to facilitate the macro- and microimpregnation processes. Finally the heating-plates were opened and the mold including the fiber package was cooled to room temperature outside the hot-press. Second approach involved the application of pressure after WGF-PBT was allowed to undergo impregnation for 10 min. The composite manufactured under pressurecontrolled (PC) condition is designated as WGF-PBT (PC). The schematic representation of this process is shown in Fig. 1b. The starting procedure was basically similar to that of DC condition. However, under the PC approach, after a defined period of time the fiber package was compressed by the preset pressure until the molding process was completed. A slight shrinkage of about 0.05 mm (ca. 1.7%) could be noticed in the pressure state. In both approaches the total molding time was fixed at 30 min. 3. Characterizations 3.1. Rheology A plate–plate rheometer (Ares, Rheometric Scientific, New Jersey, USA) was used to carry out the rheological testing. A typical procedure for the rheological measurement is as follows: The chamber was preheated to a given temperature. The dried CBT was then introduced into the platen with a diameter of 25 mm. The distance of the two

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Fig. 1. Manufacturing process for WGF-PBT composites under (a) displacement-controlled (b) pressure-controlled conditions.

plates of the fixture was adjusted to a gap height of about 1 mm. The complex viscosity and its changes as a function of time and temperature were measured. Measurement was carried out under isothermal condition at T = 190 C and the variation of rheological parameters, such as viscosity and phase angle with time, were monitored. 3.2. Differential scanning calorimetry Differential scanning calorimeter (DSC) was performed on IM-PBT and ISP-PBT on DSC 821 device, (Mettler Toledo, Giesen, Germany). Experiments were run with samples ranging from 8–10 mg under nitrogen to prevent moisture and oxidative degradation. The samples were subjected to the heating and cooling rates of 5 C/min. Crystallization and melting parameters were derived from the cooling and heating scans. Melting temperature (Tm) was considered to be the maximum of the endothermic melting peak from the heating scans, whereas crystallization temperature (Tc) was taken as the maximum of the exothermic peak of crystallization from the cooling scans. The heat of fusion (DHm) and the crystallization heat (DHc) were determined from the areas under the melting and crystallization peaks, respectively. The degree of crystallinity (Xc) of the samples was calculated by accepting 145 J/g as the heat of fusion for a 100% crystalline polymer [9]. 3.3. Density, volume fraction and void content The GF content of the composite was determined by ashing the matrix. The samples were pyrolyzed in a furnace

(Isuzu Seisakusho, Kyoto, Japan) by gradually increasing the temperature to 600 C and maintaining them at that temperature for 2 h. The void content of the samples was determined by carrying out density measurements according to ASTM D792-86, followed by burn-out tests. The density of all materials used in this study was determined using Micrometrics Accupyc 1330 Pycnometer under nitrogen gas. Microscopic observations of polished sections taken normal to the thickness of the specimens were performed to analyze the 2D geometric arrangement of fibers within the polymer matrix. Scanning electron microscopic (SEM) micrographs were taken at a 25-kV acceleration voltage at various magnifications using a Jeol JSM 5400 scanning electron microscope. Prior to the SEM observations the polished sections were mounted on aluminum stubs and sputter coated with a thin layer of gold to avoid electrical charging. 4. Testing 4.1. Mechanical Tensile and flexural tests were performed on PBT and WGF-PBT specimens in accordance with ASTM-D638 and ASTM-D790, respectively. Short beam shear tests were done according to ASTM D-2344 to estimate the apparent inter-laminar shear strength (ILSS). The ILSS was determined using a loading flexure fixture. All the mechanical testing was performed on universal testing machine (Zwick 1445, Ulm, Germany) at room temperature. The average values were derived from five parallel

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tests for each material. In order to get information on the failure, broken specimens were inspected in SEM. They have been prepared as described previously for the polished sections. 4.2. Dynamic mechanical analysis The stiffness change as a function of the temperature (T) was determined by dynamic mechanical thermal analysis (DMTA). Rectangular specimens of about 60 · 10 mm2, were subjected to load-controlled sinusoidal loading in a DMTA device (Explorer 150-N Gabo Qualimeter, Ahlden, Germany). Measurements occurred on a three-point bending mode (static and oscillating loads of 20 and 10 N, respectively; span length: 30 mm) at a frequency of 1 Hz. The composites’ viscoelastic properties i.e., the storage modulus, E 0 , the loss modulus, E00 , complex dynamic modulus, E* and the mechanical loss factor, tan d (where tan d = E00 /E 0 ) were evaluated in the temperature range of 100 to +220 C. A heating rate of 1 C/min was set for the tests. 5. Results and discussion 5.1. Differential scanning calorimetry Table 1 summarizes the melting temperatures and enthalpies associated with the first and second DSC heating, as well as the crystallization temperatures and crystallization enthalpies, deduced from the cooling regime. Fig. 2a shows typical DSC thermograms of CBT, ISPPBT and IM-PBT for the first heating at 5C/min from 0 to 250 C. A melting peak observed at 142 C is due to the melting of CBT. This is in agreement with earlier observations reported by other researchers [3,7,10]. It can be seen that CBT melts in a broad range (from about 120 to 170 C). The broadening of the shoulder may be ascribed to the different melting temperatures of the oligomers present in the CBT. According to Brunelle [11], CBT oligomers obtained via depolymerization route such as the one used in this study comprises of different proportions of oligomers indicated by slightly different melting temperatures. A small melting peak which corresponds to the melting behavior of PBT (due to the polymerization of CBT) could also be observed at 226 C. The reason for the missing crystallization (and subsequent melting peak) on the CBT trace in Fig. 2a is of kinetic origin. CBT is known to undergo concurrent crystallization with polymerization in the temperature range from 170 to 210 C. However, in the present study no evidence of any

Fig. 2. Thermograms of CBT, ISP-PBT and IM-PBT during (a) first and (b) second heating (heating rate = 5 C/min).

exothermic process linked with the in situ polymerization or crystallization was detected (c.f. Fig. 2a). In our earlier publication [12] it was reported that the heating rates play an important role in detecting the exothermic process (mostly related with the crystallization) during the first heating of CBT. At a much lower heating rate of 0.5 C/ min, the DSC thermogram showed a prominent exothermic peak at 195 C which corresponds to the crystallization of PBT formed during the ring-opening polymerization. The slower heating rate allows the polymerized CBT to crystallize and melt subsequently inside the DSC pan. So, high DSC heating rates hinder the crystallization of the PBT produced and as a consequence its melting peak is missing, too. In another study, modulated DSC (MDSC) was used to study the polymerization of a CBT as a function of the end temperature (T = 200 and 260 C) and holding time [13]. The melting behavior of the PBT was strongly affected by the polymerization mode (performed below or above the melting temperature of the PBT).

Table 1 Melting and crystallization parameters of IM-PBT and ISP-PBT First heat

CBT ISP-PBT IM-PBT

Cooling

Second heat

Tm (C)

DHm (J/g)

Xc (%)

Tc (C)

DHc (J/g)

Tm1 (C)

Tm2 (C)

DHm (J/g)

142 227.6 224.4

66.8 70.5 50.2

– 48.6 34.6

195.4 189.7 191.7

56.6 57.2 43.1

216.9 218.6 213.8

224.5 224.7 224.6

60.7 56.4 45.2

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Fig. 3 shows a typical plot for the variation of viscosity and phase angle with time for CBT at 190 C. The viscosity of CBT was about 0.02 Pa s at the start of the experiment. This value is in agreement with the reported data by Cyclics [3]. The steep rise in viscosity observed within seconds indi-

cates the fast nature of the ring-opening polymerization taking place in the melt. The abrupt change in the viscosity is accompanied by a sharp drop in the phase angle, thus indicating a dramatic change in the viscoelastic behavior of the melt. After that the viscosity remains fairly stable at a very high viscosity [12]. One of the important aspects of rheological study of CBT is to establish the impregnation and polymerization time to be used in the molding of ISP-PBT and WGFPBT. The continuous increase in the viscosity with time (c.f. Fig. 3) implies that limited time is available for the impregnation of fiber reinforcements and yet one of the main requirements to achieve good composite properties is a complete impregnation. According to the literature, direct impregnation can be applied if the matrix viscosity during the impregnation stage is very low, i.e., 1 Pa s or less [3,16]. Accordingly, an acceptable definition of the impregnation time is the time when the reactive polymer system reaches a viscosity of 1 Pa s [3,12]. The impregnation region, shown in Fig. 3, reflects those conditions whereby a good impregnation of the reinforcement with CBT could be expected. Note that it took just about 3 min to reach a viscosity of 1 Pa s at 190 C. The impregnation time of CBT could be prolonged by reducing the temperature. A detailed study on the effect of temperatures on the impregnation time of CBT is given elsewhere [12]. Since the rising of viscosity could be associated with the in situ polymerization of CBT, the rheological curve could also serve as a tool to determine the polymerization time. The fact that the phase angle of almost zero signifies the total transformation of the polymerized CBT into a solid phase via crystallization, a tangent was drawn to the phase angle curve in its late stage (see the insertion in Fig. 3). An apparent polymerization time could then be obtained from the intersection of the tangent to the curve. In order to confirm the validity of this approach, a correlation between the phase angle data obtained from the rheological studies at 190 C and the degree of conversion was established – c.f. Fig. 4. The conversion was taken equal with the in tetrahydrofuran insoluble fraction of the PBT [12]. The rapid rise

Fig. 3. Variation of viscosity and phase angle with time for CBT at 190 C.

Fig. 4. Variation of conversion and phase angle with time for CBT at 190 C.

Fig. 2a reveals the melting profiles observed during the first heating of ISP-PBT and IM-PBT. A distinct melting peak which could easily be associated with Tm of PBT was observed in both cases albeit the DHm value of IMPBT was below than that of ISP-PBT (c.f. Table 1). In the case of ISP-PBT, the existence of the melting peak indicates that complete polymerization and crystallization processes have taken place. It is interesting to note that the mean crystallinity (Xc) of ISP-PBT is higher than that of IM-PBT (cf. Table 1). Earlier studies by Brunelle et al. [11] indicated that polymers synthesized from cyclic oligomers displayed higher enthalpies of melting (60–80 J/g) hence, higher Xc as compared to the conventional counterpart with enthalpies of melting ranging from 35–50 J/g. Thus the nature of the polymerization (i.e., ring-opening polymerization versus polycondensation) plays a crucial role in determining the crystalline structure of the resulting PBT. This is a very interesting aspect which is now under investigation using X-ray diffraction techniques. Fig. 2b shows the DSC thermograms of CBT (in that case polymerized CBT), ISP-PBT and IM-PBT due to a second heating at 5 C/min. In all cases, one broad endotherm with two melting peaks could be observed (c.f. Table 1). A recent study by Righetti and Lorenzo ([14]) using a modulated DSC revealed that the melting process of PBT consists of the melting of small and defective crystals, followed by their immediate recrystallization into more stable structures and their subsequent melting. This was confirmed by our MDSC study as well [13]. In the present case a more prominent melting peak was observed at a higher temperature for IM-PBT. This is in agreement with the findings reported by other researchers [14,15]. 5.2. Rheology

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Table 2 Density, tensile and flexural data of IM-PBT and ISP-PBT

IM-PBT ISP-PBT

Density (g/cm3)

Tensile strength (MPa)

Tensile modulus (GPa)

Tensile strain (%)

Flexural strength (MPa)

Flexural modulus (GPa)

1.30 1.32

55.9 ± 3.5 58.6 ± 2.5

2.4 ± 0.2 2.3 ± 0.1

8.0 ± 1.4 2.3 ± 1.6

112.4 ± 4.2 104.2 ± 9.2

2.3 ± 0.1 2.4 ± 0.7

Table 3 Tensile, flexural and inter-laminar shear strength (ILSS) of WGF-PBT prepared under two different molding conditions

WGF-PBT(DC) WGF-PBT (PC)

Tensile modulus (GPa)

Tensile strength (MPa)

Tensile strain (%)

Flexural modulus (GPa)

Flexural strength (MPa)

Inter-laminar shear strength (MPa)

18.8 ± 0.8 20.6 ± 0.3

302 ± 5 356 ± 9

1.8 ± 0.03 1.5 ± 0.04

22.3 ± 0.1 24.5 ± 0.3

482 ± 13 578 ± 8

28.2 ± 1.5 34.3 ± 1.1

in the degree of conversion from t = 1 to t = 10 min is accompanied by a sharp drop in the phase angle from 90 to about 20. Further increase in time resulted in a marginal increase in the degree of conversion and a smaller reduction in the phase angles. This is followed by stabilization in both plots. The degree of conversion which reached ca. 95% is in agreement with literature data [3]. Based on Fig. 4 it is fairly reasonable to conclude that polymerization of CBT was completed after 30 min. Thus in this study a total molding time of 30 min was chosen in preparing ISP-PBT and WGF-PBT. Obviously, it would be interesting to investigate the effect of molding time on the microstructure–property relationships of both ISP-PBT and IM-PBT. This will be the subject of our future publication. 5.3. Mechanical properties Table 2 shows the tensile and flexural properties of IMPBT and ISP-PBT. It can be seen that both types of PBT possess similar values of strength and stiffness. However, the relatively low tensile strain of ISP-PBT indicates its brittle nature. From microstructural point of view, there are several possible reasons for this brittleness. Presence of defects, low molecular weight, high degree of crystallinity and large spherulites can substantially reduce the ductility of a polymeric material [8]. In the present case ISP-PBT shows much higher degree of crystallinity than IM-PBT – c.f. Table 1. Thus, this may account for the brittleness of ISP-PBT. The higher density of ISP-PBT is also in-line with the crystallinity data (c.f. Table 2). Table 3 summarizes the tensile properties of WGF-PBT molded under two different conditions, i.e., displacement(DC) and pressure-controlled (PC) processes. As expected, in both cases, the incorporation of WGF reinforcement resulted in a significant enhancement of both stiffness and strength, however at the expense of the tensile strain. A further improvement of the tensile strength up to 18% was achieved by WGF-PBT (PC) albeit the fact that fiber content was slightly lower than WGF-PBT (DC) – c.f. Table 4. This enhancement may be attributed to the lower amount of void content of the composites produced under pressure-control condition. Further discussion of this subject will be given later.

Typical flexural stress–deflection curves of IM-PBT, ISP-PBT and WGF-PBT are shown in Fig. 5, and the related data are summarized in Table 3. The brittleness of ISP-PBT is still obvious in flexural similar to the tensile deformation. The flexural strength of ISP-PBT is in agreement with the flexural strength of a typical injection molded PBT, which is known to be in the range of 80– 115 MPa [17]. As expected, the incorporation of WGF into ISP-PBT has resulted in a significant enhancement of both stiffness and strength. The reinforcement imparted by the fibers allows stress transfer from the matrix to the fibers. It is interesting again to note further improvement of up to 20% in the flexural strength attained by WGF-PBT produced under PC condition. The flexural strength of nearly 580 MPa recorded was quite remarkable considering the fact that the reinforcement was in the form of woven glass Table 4 Density, volume fraction of glass fibers and void content of WGF-PBT prepared under two different molding conditions

WGF-PBT(DC) WGF-PBT (PC)

Density (g/cm3)

Fiber content (vol.%)

Void content (%)

1.95 2.02

54.4 53.6

4.7 <1

Fig. 5. Typical 3 point bending stress–deflection plots for ISP-PBT, IMPBT, WGF-PBT (DC) and WGF-PBT (PC).

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Fig. 6. SEM micrograph showing the good interfacial bonding between ISP-PBT and GF in WGF-PBT (PC).

fabric. More so, if the data are to be compared with the flexural strength of 766 MPa for unidirectional fabric reinforced ISP-PBT manufactured by using the RTM technique [8]. The significant improvement of up to 5–7 times of the strength (in both tensile and flexural modes) of neat PBT already suggests a good interfacial bonding between ISP-PBT matrix and the WGF reinforcement. SEM micrograph of WGF-PBT (PC) displayed in Fig. 6 corroborates the good interfacial bonding between the constituent materials. It can be seen that the ISP-PBT matrix strongly adhered to the GF surfaces. The quality of the ISP-PBT impregnation was further investigated by polishing the cross-section of molded WGF-PBT composites and viewed under SEM. Fig. 7 shows the 2D geometric arrangement of WGF in WGF-PBT (PC). It can be observed that fibers are packed into their respective warp and weft rovings. In addition, at this analysis scale, it seems that the inter-fiber distance in the rovings remains fairly uniform. SEM micrograph in Fig. 8 reveals that a good impregnation was achieved in the WGF-PBT (PC) composites in fact. It is reasonable to suggest that the low viscosity CBT facilitates the penetration of the resin through

Fig. 7. SEM observations of 2D geometric arrangement of fibers in WGFPBT (PC) composites.

Fig. 8. SEM micrograph of polished surface showing the fiber dispersion and quality of impregnation in WGF-PBT (PC) composites.

the fabric and even more importantly, the wetting of the inner fibers within the weft or warp rovings. This perhaps explains the significant enhancement observed for the stiffness and strength of the composites (c.f. Table 3). Such good micro- and macro-impregnations were however, not observed in the case of WGF-PBT (DC). The 2D geometric arrangement of WGF shown in Fig. 9 revealed extensive microstructural defects in the form of voids owing to poor micro-impregnation. This verifies the higher void content calculated for WPC-PBT (DC) specimens (– c.f. Table 4). A further evidence of voiding in the WGF-PBT (DC) can be seen in Fig. 10. These voids could acts as stress concentration sites and leads to premature failure of the composites. It is obvious that molding under DC-control is not that effective in respect with the consolidation than PCcontrol. Consequently, the mechanical performance of WGF-PBT (DC) is inferior to WGF-PBT (PC). Thus the application of pressure is necessary to ensure good microand macro-impregnation of the GF fabric within the composite irrespective to the fact that the viscosity of CBT in

Fig. 9. SEM observations of 2D geometric arrangement of fibers in WGFPBT (DC) composites.

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Fig. 11. Variation of storage modulus, E 0 and tan d with temperature for IM-PBT, ISP-PBT and WGF-PBT (PC).

Fig. 10. SEM micrograph of polished surface showing the fiber dispersion and quality of impregnation in WGF-PBT (DC) composites.

the impregnation stage is very low (– c.f. Fig. 3). The application of pressure facilitates the flow of the low viscosity CBT to penetrate into the warp and weft rovings and ensuring the complete wetting of their constituting fibers. Table 3 shows the apparent inter-laminar shear strength (ILSS) values obtained from the short beam shear test for WGF-PBT (DC) and WGF-PBT (PC). The ILSS values obtained in the present study (28.2–34.4 MPa) are within the range of ILSS reported by other workers for woven glass fabric composites. Pavlidou et al. [18] quoted ILSS value of 15 MPa for plain-woven fabric/unsaturated polyester composites. Another study by Kim et al. [19] reported ILSS values ranging from about 35–50 MPa for plainwoven glass fiber reinforced vinylester composites treated with different types of silane coupling agents. Karger-Kocsis reported interfacial shear strength between 7 and 20 MPa for GF fabric reinforced PP composites produced from commingled yarn [20]. The slightly higher ILSS value of WGF-PBT (PC) as compared to that of WGF-PBT (DC) clearly support the mechanical data obtained from both tensile and flexural tests. Higher ILSS value of WGF-PBT (PC) implies better efficiency of stress-transfer from the PBT to the WGF reinforcements which results in higher strength as compared to that of WGF-PBT (DC). In the past, several studies have shown that increasing the level of fiber-matrix bonding (by tailoring of the interfacial region with different sizings) can yield significant improvements in the mechanical properties of composites [21]. This will be the subject of our future publication. The temperature dependence of the storage modulus, E 0 and tan d and its response to the presence of WGF reinforcement is illustrated for ISP-PBT in Fig. 11. The dynamic mechanical behavior of both ISP-PBT and IMPBT are also included for comparison. As discussed in the literature, the glass transition temperature (Tg) of crystalline polymers may increase or decrease with crystallinity depending on the nature and the role of crystalline phase in

the glass transition process. The tan d peaks of ISP-PBT and IM-PBT which are located at 60 C and 64 C, respectively, indicate that the molecular segments of the related Tg relaxations may be different. The effect of crystallinity is more obvious by comparing the height of the tan d peak between the two types of PBT. The depression observed in the tan d of ISP-PBT as compared to IM-PBT may be attributed to the higher degree of crystallinity of the former. The tan d maximum of WGF-PBT composite is located at about the same temperature as that recorded for the ISP-PBT. A further depression in the intensity of tan d peak of ISP-PBT could be observed with the incorporation of WGF. This effect may be attributed to a reduction of the mobility of the macromolecular chains in the vicinity of the glass surface due to good interfacial interaction between ISP-PBT and WGF. Another effect of the reinforcement is the flattening and broadening of the glass transition region. Again, this can be assigned to the formation of an interphase between the GF and bulk matrix [22]. The properties of the interphase usually differ from that of the bulk owing to adsorption, crystallization phenomena (e.g. [22,23]). The increase of E 0 over the analyzed temperature range can be attributed to the usual reinforcement effect. Good stiffening effect was achieved particularly at high temperature. Such behavior may be explained by the fact that while modulus of the GF is not affected by changes in temperature, the PBT matrix undergoes a sharp decrease in modulus with increasing temperature above Tg. 6. Conclusions Isothermal processing of PBT and WGF-PBT via in situ polymerization of CBT has been successfully performed using compression molding technique. Both IM-PBT and ISP-PBT seem to display similar strength and stiffness values under tensile and flexural testing. The relatively higher degree of crystallinity of ISP-PBT is believed to be responsible for the brittle failure of this material. The application of pressure during the molding of WGF-PBT resulted in a further improvement in the stiffness and strength of the composites. SEM examinations not only revealed good

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interfacial bonding between ISP-PBT matrix and WGF (as also confirmed by the ILSS value from short beam shear test) but also uniform fiber dispersion in the matrix. Thus, it can be concluded that the low viscosity CBT does not only facilitate the penetration of the resin through the fabric but more importantly also the wetting of the inner fibers within the rovings of the glass fabric. Acknowledgements Z.A. Mohd Ishak is thankful to the Alexander von Humboldt Foundation and Universiti Sains Malaysia for his Georg Forster Research Fellowship and for granting his sabbatical leave, respectively, at the IVW. The fellowship of the German Academic Exchange Service (Deutscher Akademischer Austauschdienst) to Y.W. Leong is gratefully acknowledged. Part of this work was made in the framework of a BMBF Project (Pro-PBT). References [1] Lou AY, Murtha TP, O’Connor JE, Brady DG. Continuous-fibre thermoplastic composites. In: Carlsson LA, editor. Thermoplastic composites materials. Composite materials series 7. Amsterdam: Elsevier Science Publishers B.V.; 1991. p. 167–204. [2] Hou M, Ye L, Mai YW. Manufacturing process and mechanical properties of thermoplastic composite components. J Mater Process Tech 1997;63:334–8. [3] Eder RHJ. Cyclic thermoplastics-properties and processing.In: Neitzel M, editor.‘‘Flu¨ssigimpra¨gnierung mit Duro- und Thermoplasten’’ IVW Band 25, Universita¨t Kaiserslautern, 2001, p. 33–43. [4] Tripathy AR, Burgaz E, Kukureka SN, MacKnight WJ. Poly(butylene terephthalate) nanocomposites prepared by in situ polymerization. Macromolecules 2003;38:8593–5. [5] Liu Y, Wang YF, Gerasimov TG, Heffner KH, Harmon JP. Thermal analysis of novel underfill materials with optimum processing characteristics. J Appl Polym Sci 2005;98:1300–7. [6] Tripathy AR, Chen W, Kukureka SN, MacKnight WJ. Novel poly(butylene terephthalate)/poly(vinyl butyral) blends prepared by in situ polymerization of cyclic poly(butylene terephthalate) oligomers. Polymer 2003;44:1835–42. [7] Parton H, Verpoest I. In situ polymerization of thermoplastic composites based on cyclic oligomers. Polym Comp 2005;26: 60–65.

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