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polypropylene composites
Composites: Part B 54 (2013) 415–421
Contents lists available at SciVerse ScienceDirect
Composites: Part B journal homepage: www.elsevier.com/locate/compositesb
Flexural and impact response of woven glass fiber fabric/polypropylene composites P. Russo a,⇑, D. Acierno b, G. Simeoli b, S. Iannace c, L. Sorrentino c a
Institute of Chemistry and Technology of Polymers, National Council of Research, Via Campi Flegrei 34, 80078 Pozzuoli, Naples, Italy Department of Chemical, and Materials Engineering and Industrial Production, University of Naples Federico II, P.le V. Tecchio 80, 80125 Naples, Italy c Institute of Composite and Biomedical Materials, National Council of Research, P.le E. Fermi 1, 80055 Portici, Naples, Italy b
a r t i c l e
i n f o
Article history: Received 3 January 2013 Accepted 9 June 2013 Available online 20 June 2013 Keywords: A. Polymer–matrix composites A. Laminates A. Glass fibers B. Impact behavior B. Mechanical properties
a b s t r a c t Impact tests with a falling dart and flexural measurements were carried out on polypropylene based laminates reinforced with glass fibers fabrics. Research has shown that the strong fiber/matrix interface obtained through the use of a compatibilizer increased the mechanical performance of such composite systems. The improved adhesion between fibers and matrix weakly affects the flexural modulus but strongly influences the ultimate properties of the investigated woven fabric composites. In fact, bending tests have shown a clear improvement in the flexural strength for the compatibilized systems, in particular when a high viscosity/high crystallinity polypropylene was used. On the contrary, the low velocity impact tests indicated an opposite dependence on the interface strength, and higher energy absorption in not compatibilized composites was detected. This result has been explained in terms of failure mechanisms at the fiber/matrix interface, which are able to dissipate large amounts of energy through friction phenomena. Pull-out of fibers from the polypropylene matrices have been evidenced by the morphological analysis of fracture surfaces after failure and takes place before the fibers breakage, as confirmed by the evaluation of the ductility index. Ó 2013 Elsevier Ltd. All rights reserved.
1. Introduction Thermoplastic composites have gained an increasing interest since the mid-80s. Due to their clear advantages over thermoset composites in terms of short processing times [1–3], improved fracture toughness, high damage tolerance and good resistance to micro cracking [4], potential recyclability and possibility to be reshaped or remolded at high temperatures, thermoplastic composites appear to have relevant perspectives in many industrial fields [5–8]. In this frame, with special regard to structural applications (i.e. transportation and construction fields), woven fabric composites have been recognized more competitive than unidirectional composites due to their good stability and deformation characteristics. Among the thermoplastic composites, the system polypropylene–glass fiber has been particularly analyzed [9–12]. As well known, glass fibers offer, in addition to a low cost, a high tensile strength, high chemical resistance and fairly good mechanical properties. However, such reinforcement has an elastic modulus lower than that of carbon fibers, a low resistance to fatigue and it usually shows a poor adhesion at the fiber–matrix interface [13,14]. This drawback, which induces poor mechanical properties, may be overcome by physical treatments [15,16] or by using ⇑ Corresponding author. Tel.: +39 0817682268. E-mail address: [email protected] (P. Russo). 1359-8368/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.compositesb.2013.06.016
suitable coupling agents like maleic anhydride to enhance the wetting effect of the hosting matrix on fibers [17,18]. Pull-out tests, for example, have demonstrated that the resistance at the fiber– matrix interface is more efficiently influenced by the presence of a compatibilizing agent rather than by using sizing on the fibers [13]. Another performance, usually decisive for estimating the effective applications of composite materials and traditionally considered dependent on the level of adhesion at the fiber–matrix interface, is the impact resistance. At this regard, a growing concern to design engineers is focused on the so-called low-velocity (i.e. <6 m/s) impact behavior since it simulates real events occurring during fabrication, maintenance and operations with such composite systems [19–21]. Damage resulting from low-velocity impacts is usually not visible from the impacted side of the structure; however, significant delamination and back face damage may be present. In particular, low-velocity impact is considered most threatening to composite structures, because the entity of the damage might easily be undetected during routine visual inspection and thus damaged parts may improperly be put into service. Padaki et al. [20], focusing the attention on textile reinforced composites, have reviewed the influences of instrument parameters and composite material variables on their complex impact behavior. Overall, starting from well established assumptions as the typical high complexity of the distribution of stress as a
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function of time and space for the reference materials, authors have emphasized the simultaneous occurrence of several interacting damage modes for composites subjected to impact events. Belingardi and Vadori [21] analyzed the behavior of composite plates made by carbon–epoxy laminates subjected to low velocity impact loading. In particular, the influence of stacking sequence on some impact parameters as the saturation impact energy and the damage degree has been investigated. Results showed that the first parameter increases with the laminate thickness at a power greater than unity whereas the degree of damage indicates a general linear trend. In this work, the behavior of polypropylene based laminates subjected to flexural and low velocity impact tests has been investigated. This research investigated the effects of the use of a compatibilizer on the mechanical behavior of laminates based on two grades of polypropylene, differing for their viscosity and crystallinity, as thermoplastic matrix and reinforced with glass fiber fabrics, with a 0°/90° configuration. Mechanical results, interpreted also on the basis of morphological observations and collected by scanning electron microscopy of fracture sections, have emphasized the strong influence of pull-out phenomena on the energy absorption during impact tests.
2. Experimental 2.1. Materials Two different grades of polypropylene, commercially available as EP348U and MA712 and supplied by Lyondell Basell Industries and Unipetrol and herein coded as M–LV and M–HV, respectively, were used as matrices for the preparation of the fiber reinforced composites. Some physical properties of these resins, taken from their datasheets, are summarized in Table 1. Polypropylene grafted with maleic anhydride (PP-g-MA) commercialized under the trade
Table 1 Mechanical properties of matrices from manufacturer’s datasheets. Properties
Method
Unit
M–LV
M–HV
Density Melt flow index (MFI) Tensile modulus Tensile stress at yield Tensile strain at break Charpy notched impact strength (+23 °C)
ISO ISO ISO ISO ISO ISO
g/cm3 g/10 min MPa MPa % kJ/m2
0.90 70 1200 24 30 5.5
0.90 12 1550 29 50 10
1183 1133 527-1/-2 527-1/-2 527-1/-2 179/1eA
name Polybond 3200 (MFI 115 g/10 min, 1 wt% maleic anhydride; from Chemtura, Philadelphia – PA, USA) was used as compatibilizer to improve the fiber/matrix interface. Compatibilized blends have been coded by adding ‘‘-PB’’ to the matrix code. Finally, a plain weave type woven glass fabric (E-type glass fibers having density of 2.54 g/cm3) with a specific mass of 204 g/m2 have been used as the reinforcement. The fabric was functionalized by amino silane groups. Composites have been coded by prefixing ‘‘C-’’ to the matrix code. 2.2. Sample preparation Films of neat or pre-blended polypropylene with 2 wt% of compatibilizer, with a thickness equal to 35–40 lm, were prepared by using film blowing extrusion line model Teach-Line E 20 T from Collin Gmbh (Ebersberg, Germany). Composite laminates were obtained alternating layers of polypropylene films and glass fiber fabrics by the film-stacking technique using a compression molding machine (model P300P, Collin Gmbh, Ebersberg, Germany) according to a pre-optimized molding cycle (see Fig. 1). With the above cited approach, plaques consisting of 8 balanced fabric layers 0°/90°, symmetrically arranged with respect to the middle plane of the laminate ([(0/90)4]s configuration), were produced with a target thickness of 1.30 mm and a glass fiber content of 50% by volume. The volume percentages of fiber and matrix were evaluated according to the ASTM D 3171-99 and ASTM D792. 2.3. Characterization techniques Calorimetric analysis was performed by using a DSC, Model 1 from Mettler Toledo Inc. (Delaware, USA). Tests were conducted on 8–10 mg samples from room temperature to 220 °C at the heating rate of 10 °C/min. Thermograms were used to evaluate main thermal parameters as the melting temperature and enthalpy and the degree of crystallization of all investigated materials, evaluated using the following equation:
DH f X c ¼ 100 DH0f
!
in which DH0f ¼ 207 J=g is the theoretical value of the melting enthalpy for the completely crystalline polypropylene. Flexural properties were carried out by means of a three point bending configuration set according to the ASTM D 790-10, using an INSTRON mod. 3360 dynamometer (Akron – OH, USA) equipped with a 5 kN load cell. The span was changed according to the actual
Fig. 1. Processing conditions used to prepare LV (A) and HV (B) based composites.
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measured thickness for each sample in order to keep the span-todepth ratio equal to 40. Flexural modulus (EF) and flexural strength (rF) values were averaged on 5 samples for each investigated material. Low-velocity impact tests were conducted using an instrumented drop-weight impact testing machine (model Fractovis Plus from CEAST – Italy) equipped with a hemispherical tip (diameter 12.7 mm). All tests were performed putting the sample on a stainless steel annular ring (internal diameter 40 mm, outer diameter 60 mm) and using an impact energy of 31 J (impact parameters: mass = 6.926 kg, velocity = 2.98 m/s). A minimum of 4 samples for each composition, measuring 80 80 mm2 and cut from the prepared laminates, were tested and their mean values and variance were calculated. Optical analysis in reflection mode was carried out by using a microscope (BX51 from Olympus, Japan) to investigate the fibers impregnation in the composite and the distribution of particles (detectable with this optical technique) dispersed in the composite matrices. Fractured surfaces were acquired by a field emission scanning electron microscope, model QUANTA-200FEG from FEI (Eindhoven, The Netherlands). The examined surfaces were coated with a thin layer of a gold-palladium alloy prior to SEM analysis. All glass–fiber composite samples were prepared by polishing the observation surfaces with wet sandpaper and then with a very fine polishing paste.
Table 2 Relevant thermal properties of all matrix compositions. Matrix
Tm,onset (°C)
Tm,peak (°C)
DHm (J/g)
Xc (%)
M–LV M–HV M–LV–PB M–HV–PB
155.8 156.2 154.7 155.7
167.3 169.3 168.1 169.8
69.7 81.9 71.4 78.5
33.7 39.6 34.5 37.9
fiber volume content with respect to the not compatibilized ones (C–LV). On the contrary, C–HV–PB samples exhibited a higher fiber content with respect to the not compatibilized composite (C–HV). This will affect the flexural response of the composite, as shown in the following. Composites were also observed to check the presence of voids within fiber bundles and in the polymer matrix. Impregnation of fabrics resulted to be very good despite of their very different melt flow index, as shown in Fig. 3: no voids were detected by several optical acquisitions in any laminate and all bundles appeared to be well impregnated by the hosting matrix. Tests were performed to measure the percentage of voids with digestion methodology according to the ASTM D3171-99 standard but results were all below the 2% limit of acceptance for the composite (Table 4). Furthermore, the fiber volume content was evaluated according to two different standard (namely ASTM D3171-99 and ASTM D792), evidencing very good accordance between them.
3. Results and discussion 3.2. Flexural properties 3.1. Thermal properties DSC heating thermograms of all matrix compositions are shown in Fig. 2, and relevant thermal properties (melting temperature onset Tm,onset, melting temperature peak Tm,peak, melting enthalpy DHm and relative crystallinity Xc) and summarized in Table 2. All materials showed a single endothermic peak, whose position is only slightly influenced by the inclusion of the compatibilizer, which reduces the melting temperature onset and increases the melting temperature peak. The degree of crystallinity of LV polypropylene was lower than that of HV polypropylene. The addition of compatibilizer marginally affected Xc even if it had an opposite behavior on LV and HV matrices, being increased and decreased, respectively. Density measurements on composites are shown in Table 3. LV compatibilized laminates (C–LV–PB) were characterized by a lower
Fig. 2. DSC heating thermograms of all matrix compositions.
The stress–strain curves from the flexural characterization of matrices are shown in Fig. 4. Neat samples did not break at 5% of strain, and only flexural modulus and strength were evaluated (Table 4). Neat polymers exhibited a higher modulus with respect to compatibilized ones, which acted as a plasticizer. Both flexural modulus and strength of high viscosity polypropylene matrix (M–HV) were the highest among the tested matrices, but its blending with the compatibilizer significantly decreased the bending stiffness. Bending tests performed on composite led to the breaking of samples for the failure of the upper side (glass fiber breakage in compression) as shown in Fig. 5, where broken bundles are evidenced. C–LV samples exhibited lower flexural modulus with respect to the C–HV laminate, because of its lower amount of fiber content and lower matrix stiffness. On the contrary, the flexural strength was significantly higher. The use of compatibilized matrices had positive effect on both flexural modulus and strength (Fig. 6). The improvement is clear in C–LV–PB samples with respect to C–LV due to their matching compositions, having the same fiber volume content. More tricky was the analysis of the C–HV and C– HV–PB data, because the latter was characterized by a slightly higher fiber content. The use of compatibilizer resulted in an improvement of the flexural strength, but it was not clear if the flexural modulus was improved for the better compatibility of fibers with the matrix or for the higher amount of fibers. As also found in the literature [9], the compatibilizer marginally improves the elastic modulus, but it leads to a more pronounced increase of both flexural strength and strain at yield. These results come from the enhanced capability, in compatibilized systems, of load transfer between the matrix and the woven fabric through the interface. The marked damaging of the upper surface of composites after the peak during flexural tests, as exemplified in Fig. 5, comes from the compressive failure of glass fibers, located within the upper layer, occurring at a lower stress in non compatibilized composites with respect to compatibilized ones. After the peak stress, a lowering of the stress-strain curve was detected, which is due to the propagation of both interlaminar and intralaminar damaging of
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Table 3 Density and fiber volume measurements on composites. Matrix
Thickness (mm)
ASTM D 3171-99 3
C–LV C–LV–PB C–HV C–HV–PB
1.30 ± 0.04 1.32 ± 0.07 1.24 ± 0.04 1.20 ± 0.01
ASTM D792
Density (g/cm )
Vf (%)
Void content (%)
Density (g/cm3)
Vf (%)
1.69 ± 0.01 1.69 ± 0.00 1.73 ± 0.02 1.75 ± 0.00
48.99 ± 0.52 48.48 ± 0.32 51.18 ± 1.04 52.39 ± 0.24
1.16 ± 0.17 0.90 ± 0.35 0.46 ± 0.11 0.79 ± 0.49
1.70 ± 0.01 1.69 ± 0.00 1.74 ± 0.02 1.76 ± 0.01
48.46 ± 0.478 47.92 ± 0.150 50.87 ± 1.114 52.17 ± 0.445
Fig. 3. Impregnation of fiber bundles in the different laminate compositions: (A) C–LV, (B) C–LV–PB, (C) C–HV and (D) C–HV–PB.
Table 4 Relevant data evaluated from the flexural tests performed on investigated composites. Sample
Flexural modulus (MPa)
Flexural strength (MPa)
Strain at yield (%)
M–LV M–LV–PB M–HV M–HV–PB C–LV C–LV–PB C–HV C–HV–PB
1149 ± 40 1136 ± 42 1409 ± 92 1329 ± 90 15708 ± 234 16197 ± 405 16276 ± 530 16808 ± 434
26.1 ± 0.5 25.9 ± 0.6 31.7 ± 0.8 30.5 ± 0.8 143.1 ± 3.2 172.9 ± 6.1 129.6 ± 8.2 179.8 ± 9.6
1.43 ± 0.10 1.74 ± 0.09 1.22 ± 0.13 1.62 ± 0.15
Fig. 5. Laminate failure after flexural tests. Fiber bundle breakages are evidenced in the upper side of the sample.
3.3. Low-velocity impact tests
Fig. 4. Stress–strain curves of flexural tests on matrices.
glass fibers and fiber/matrix interfaces, until the failure reaches the lower composite surface.
Low velocity impact tests have been performed by using an impact energy value higher than that bearable by any of the prepared composites, in order to measure the impact strength of the different configurations and to evaluate the force versus deflection and energy versus time curves up to the sample penetration. The impact characterization has shown a different behavior of the structural response with respect to the bending test. In fact, a lower peak force, even if coupled to a higher slope of the force/deflection curve has been exhibited by compatibilized samples with respect to not compatibilized laminates. This has been explained as a result of the increased flexural modulus induced by the presence of compatibilizer. The best performance in terms of dissipated energy has been shown by the C–HV samples, which also reported the highest peak force (Fig. 7). On the contrary, C–HV–PB samples
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the normal forces on the fibers and the friction during fiber slipping justifying the increased energy dissipated until the force peak is reached during impacts on not compatibilized HV samples with respect to not compatibilized LV ones. 3.4. Morphological analysis
Fig. 6. Flexural stress–strain curves for the prepared composites.
exhibited the lowest impact energy absorption after impact, probably as a result of the strong fiber/interface interaction allowed by the compatibilizer. The absorbed energies, evaluated from curves of Fig. 7, are reported as a function of time in Fig. 8. Composites prepared without the compatibilizer absorbed more energy (from 20% to 60%) than their compatibilized ones, as a result of higher load bearing and larger deformations allowed by fibers slipping (Table 5). The strong interaction between fibers and matrix hinders slipping of fibers (fiber pull-out) and it is not able to stop the propagation of the damage. On the contrary, in not compatibilized systems the crack propagation is retarded by the early slipping of the fiber/matrix interface, which results in a post-posed failure of the composite as a whole and in higher deflections and force peaks. Ultimately, the overall laminate toughness of not compatibilized samples resulted to be significantly increased by the fiber pull-out phenomena which is also believed to allow a rearrangement of the fabric to better bear the load, leading to more resilient structures and increased dissipated energies. A slight contribution to the different impact behavior between not compatibilized LV and HV samples could be attributed to their difference in crystallinity. In fact, presumably the higher degree of crystallinity of HV samples with respect to LV ones could lead to higher densification of the surrounding amorphous regions and, consequently, might induce a tightening of the matrix onto the glass fibers. These effects increase
Fig. 7. Force versus deflection curves of composites from the low velocity impact tests.
The analysis performed on the force/deflection curves has been confirmed by SEM micrographs of the fracture surfaces of impacted samples. The improved adhesion of matrix on fibers is evident in Fig. 9B and D, respectively for LV and HV based composites. In laminates produced without the compatibilizer, the fibers exhibit very smooth surfaces, confirming their slippage in the matrix as shown in Fig. 9A and C for LV and HV based composites, respectively. When the high viscosity PP was used, surfaces were very weakly impregnated, as a result of the poorer fiber/matrix interaction. SEM analysis also showed that fracture surfaces of compatibilized composites exhibited a large number of short fibers protruding from the fractured surfaces, unlike not compatibilized ones. This could be an evidence that crack propagation through fibers during impact was very fast and straight through the laminate in the compatibilized composites. On the contrary, in not compatibilized samples the weak interactions between the two phases allowed the pull-out mechanism during impact. The fibers pulling out from the matrix can be considered the main reason for the improved amount of absorbed energy during the impact, because the contribution of the friction (dissipative mechanism allowed by the relative motion of fiber/matrix surfaces) between the sliding fibers and the matrix is not present in composites with compatibilized matrices. This phenomenon resulted more pronounced in the case of high viscosity PP composites (C–HV). It is worth to note that damage zone after impact, depicted in Fig. 10, presents cross shaped cracks in both compatibilized samples, as usually occurs in composites with strong adhesion between fibers and matrix, while in not compatibilized samples damages and cracks are evenly distributed around the impacted hole, not presenting preferential directions of crack propagation along the main directions of the glass fiber fabric. 3.5. Ductility index Further considerations on the impact behavior of the prepared composite systems can be obtained comparing their ductility index (Table 5), defined as the ratio between the energy absorbed after the force peak (Etot Epeak) and the energy absorbed up to the peak
Fig. 8. Absorbed energy versus time curves of composites from the low velocity impact tests.
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Table 5 Data from low velocity impact characterization. Sample
C–LV C–LV–PB C–HV C–HV–PB
Properties at maximum load Load (N)
Deflection (mm)
Epeak
Etot Epeak
Ductility index (DI)
2863 ± 69 2670 ± 52 3228 ± 36 2487 ± 56
5.94 ± 0.14 5.28 ± 0.21 6.81 ± 0.26 5.42 ± 0.26
7.74 ± 0.42 6.36 ± 0.35 9.86 ± 0.38 5.95 ± 0.32
8.71 ± 0.13 8.64 ± 0.71 9.82 ± 1.47 9.53 ± 0.17
1.13 ± 0.07 1.36 ± 0.12 1.00 ± 0.16 1.60 ± 0.06
energy dissipated by mechanisms not involving the breaking of the reinforcing fibers, while after the force peak large fibers breaking occurs. Results demonstrated that the presence of the compatibilizer increases the ductility index by decreasing the area under the force-deflection curve before its maximum with respect to the area under the same curve at high deflections. This behavior seemed to be reversed for not compatibilized systems as a result of the increased energy dissipated before large fiber breaking. In fact, typical dissipative mechanisms, involved in these latter samples and related to the slipping of fibers within the matrix, occur only during the early stage of the impact before fibers failure. In other words, a higher amount of energy is absorbed in the fiber’s elastic region in not compatibilized laminates with respect to compatibilized ones, thanks to the occurrence of pullout phenomena, which in turn seem to occur during the early stage of the impact, before fibers or matrix failures.
4. Conclusions
Fig. 9. Fracture surfaces of impacted samples (A) C–LV, (B) C–LV–PB, (C) C–HV and (D) C–HV–PB.
Fig. 10. Optical pictures of perforated samples (A) C–LV, (B) C–LV–PB, (C) C–HV and (D) C–HV–PB.
force (Epeak), calculated as the integral of the force with respect to the deflection. The first part of the curve represents the amount of
Composite laminates have been produced by using two types of polypropylene matrices, differing for their melt flow index, and a glass fiber fabric. Their static and low velocity impact properties have been evaluated and compared to those exhibited by composite laminates produced by adding a compatibilizing agent to the matrices. The higher amount of crystalline phase, detected by thermal analysis, was responsible for the higher flexural modulus exhibited by the high viscosity PP with respect to the low viscosity one. This consideration also affects laminates behavior and it could have an important contribution on their improved impact behavior. The use of the compatibilizer in both neat matrices and composites resulted in improved static structural properties, namely flexural modulus and flexural strength. While the flexural modulus was only marginally improved by the presence of the compatibilizer (3% increase), the flexural strength was strongly affected, up to an increase of the flexural strength in excess of 30%, in reinforced laminates. The reduced flexural strength in not compatibilized systems has been related to the low capability to transfer the load from the polymeric matrix to the reinforcing fabric. The low velocity impact tests, performed by perforating samples, have shown that the interface between fibers and matrix has a primary role in increasing the impact strength. Force vs deflection curves show that compatibilized systems exhibit higher slopes, as a consequence of their higher stiffness, but incur failure at lower forces with respect to not compatibilized systems. The reduced interface strength in the latter between fibers and matrix (also confirmed by the SEM analysis of composites’ fracture surfaces) was responsible for the occurrence of fibers pull-out phenomena, which dissipate high amount of energy through the friction consequent to sliding mechanisms. The ductility index, evaluated from the impact curves, established that the dissipative mechanisms through relative motion between fibers and matrix were the only reason for the improved impact strength in not compatibilized composites, and that they occurred before fibers failure.
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Contents lists available at SciVerse ScienceDirect
Composites: Part B journal homepage: www.elsevier.com/locate/compositesb
Flexural and impact response of woven glass fiber fabric/polypropylene composites P. Russo a,⇑, D. Acierno b, G. Simeoli b, S. Iannace c, L. Sorrentino c a
Institute of Chemistry and Technology of Polymers, National Council of Research, Via Campi Flegrei 34, 80078 Pozzuoli, Naples, Italy Department of Chemical, and Materials Engineering and Industrial Production, University of Naples Federico II, P.le V. Tecchio 80, 80125 Naples, Italy c Institute of Composite and Biomedical Materials, National Council of Research, P.le E. Fermi 1, 80055 Portici, Naples, Italy b
a r t i c l e
i n f o
Article history: Received 3 January 2013 Accepted 9 June 2013 Available online 20 June 2013 Keywords: A. Polymer–matrix composites A. Laminates A. Glass fibers B. Impact behavior B. Mechanical properties
a b s t r a c t Impact tests with a falling dart and flexural measurements were carried out on polypropylene based laminates reinforced with glass fibers fabrics. Research has shown that the strong fiber/matrix interface obtained through the use of a compatibilizer increased the mechanical performance of such composite systems. The improved adhesion between fibers and matrix weakly affects the flexural modulus but strongly influences the ultimate properties of the investigated woven fabric composites. In fact, bending tests have shown a clear improvement in the flexural strength for the compatibilized systems, in particular when a high viscosity/high crystallinity polypropylene was used. On the contrary, the low velocity impact tests indicated an opposite dependence on the interface strength, and higher energy absorption in not compatibilized composites was detected. This result has been explained in terms of failure mechanisms at the fiber/matrix interface, which are able to dissipate large amounts of energy through friction phenomena. Pull-out of fibers from the polypropylene matrices have been evidenced by the morphological analysis of fracture surfaces after failure and takes place before the fibers breakage, as confirmed by the evaluation of the ductility index. Ó 2013 Elsevier Ltd. All rights reserved.
1. Introduction Thermoplastic composites have gained an increasing interest since the mid-80s. Due to their clear advantages over thermoset composites in terms of short processing times [1–3], improved fracture toughness, high damage tolerance and good resistance to micro cracking [4], potential recyclability and possibility to be reshaped or remolded at high temperatures, thermoplastic composites appear to have relevant perspectives in many industrial fields [5–8]. In this frame, with special regard to structural applications (i.e. transportation and construction fields), woven fabric composites have been recognized more competitive than unidirectional composites due to their good stability and deformation characteristics. Among the thermoplastic composites, the system polypropylene–glass fiber has been particularly analyzed [9–12]. As well known, glass fibers offer, in addition to a low cost, a high tensile strength, high chemical resistance and fairly good mechanical properties. However, such reinforcement has an elastic modulus lower than that of carbon fibers, a low resistance to fatigue and it usually shows a poor adhesion at the fiber–matrix interface [13,14]. This drawback, which induces poor mechanical properties, may be overcome by physical treatments [15,16] or by using ⇑ Corresponding author. Tel.: +39 0817682268. E-mail address: [email protected] (P. Russo). 1359-8368/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.compositesb.2013.06.016
suitable coupling agents like maleic anhydride to enhance the wetting effect of the hosting matrix on fibers [17,18]. Pull-out tests, for example, have demonstrated that the resistance at the fiber– matrix interface is more efficiently influenced by the presence of a compatibilizing agent rather than by using sizing on the fibers [13]. Another performance, usually decisive for estimating the effective applications of composite materials and traditionally considered dependent on the level of adhesion at the fiber–matrix interface, is the impact resistance. At this regard, a growing concern to design engineers is focused on the so-called low-velocity (i.e. <6 m/s) impact behavior since it simulates real events occurring during fabrication, maintenance and operations with such composite systems [19–21]. Damage resulting from low-velocity impacts is usually not visible from the impacted side of the structure; however, significant delamination and back face damage may be present. In particular, low-velocity impact is considered most threatening to composite structures, because the entity of the damage might easily be undetected during routine visual inspection and thus damaged parts may improperly be put into service. Padaki et al. [20], focusing the attention on textile reinforced composites, have reviewed the influences of instrument parameters and composite material variables on their complex impact behavior. Overall, starting from well established assumptions as the typical high complexity of the distribution of stress as a
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function of time and space for the reference materials, authors have emphasized the simultaneous occurrence of several interacting damage modes for composites subjected to impact events. Belingardi and Vadori [21] analyzed the behavior of composite plates made by carbon–epoxy laminates subjected to low velocity impact loading. In particular, the influence of stacking sequence on some impact parameters as the saturation impact energy and the damage degree has been investigated. Results showed that the first parameter increases with the laminate thickness at a power greater than unity whereas the degree of damage indicates a general linear trend. In this work, the behavior of polypropylene based laminates subjected to flexural and low velocity impact tests has been investigated. This research investigated the effects of the use of a compatibilizer on the mechanical behavior of laminates based on two grades of polypropylene, differing for their viscosity and crystallinity, as thermoplastic matrix and reinforced with glass fiber fabrics, with a 0°/90° configuration. Mechanical results, interpreted also on the basis of morphological observations and collected by scanning electron microscopy of fracture sections, have emphasized the strong influence of pull-out phenomena on the energy absorption during impact tests.
2. Experimental 2.1. Materials Two different grades of polypropylene, commercially available as EP348U and MA712 and supplied by Lyondell Basell Industries and Unipetrol and herein coded as M–LV and M–HV, respectively, were used as matrices for the preparation of the fiber reinforced composites. Some physical properties of these resins, taken from their datasheets, are summarized in Table 1. Polypropylene grafted with maleic anhydride (PP-g-MA) commercialized under the trade
Table 1 Mechanical properties of matrices from manufacturer’s datasheets. Properties
Method
Unit
M–LV
M–HV
Density Melt flow index (MFI) Tensile modulus Tensile stress at yield Tensile strain at break Charpy notched impact strength (+23 °C)
ISO ISO ISO ISO ISO ISO
g/cm3 g/10 min MPa MPa % kJ/m2
0.90 70 1200 24 30 5.5
0.90 12 1550 29 50 10
1183 1133 527-1/-2 527-1/-2 527-1/-2 179/1eA
name Polybond 3200 (MFI 115 g/10 min, 1 wt% maleic anhydride; from Chemtura, Philadelphia – PA, USA) was used as compatibilizer to improve the fiber/matrix interface. Compatibilized blends have been coded by adding ‘‘-PB’’ to the matrix code. Finally, a plain weave type woven glass fabric (E-type glass fibers having density of 2.54 g/cm3) with a specific mass of 204 g/m2 have been used as the reinforcement. The fabric was functionalized by amino silane groups. Composites have been coded by prefixing ‘‘C-’’ to the matrix code. 2.2. Sample preparation Films of neat or pre-blended polypropylene with 2 wt% of compatibilizer, with a thickness equal to 35–40 lm, were prepared by using film blowing extrusion line model Teach-Line E 20 T from Collin Gmbh (Ebersberg, Germany). Composite laminates were obtained alternating layers of polypropylene films and glass fiber fabrics by the film-stacking technique using a compression molding machine (model P300P, Collin Gmbh, Ebersberg, Germany) according to a pre-optimized molding cycle (see Fig. 1). With the above cited approach, plaques consisting of 8 balanced fabric layers 0°/90°, symmetrically arranged with respect to the middle plane of the laminate ([(0/90)4]s configuration), were produced with a target thickness of 1.30 mm and a glass fiber content of 50% by volume. The volume percentages of fiber and matrix were evaluated according to the ASTM D 3171-99 and ASTM D792. 2.3. Characterization techniques Calorimetric analysis was performed by using a DSC, Model 1 from Mettler Toledo Inc. (Delaware, USA). Tests were conducted on 8–10 mg samples from room temperature to 220 °C at the heating rate of 10 °C/min. Thermograms were used to evaluate main thermal parameters as the melting temperature and enthalpy and the degree of crystallization of all investigated materials, evaluated using the following equation:
DH f X c ¼ 100 DH0f
!
in which DH0f ¼ 207 J=g is the theoretical value of the melting enthalpy for the completely crystalline polypropylene. Flexural properties were carried out by means of a three point bending configuration set according to the ASTM D 790-10, using an INSTRON mod. 3360 dynamometer (Akron – OH, USA) equipped with a 5 kN load cell. The span was changed according to the actual
Fig. 1. Processing conditions used to prepare LV (A) and HV (B) based composites.
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measured thickness for each sample in order to keep the span-todepth ratio equal to 40. Flexural modulus (EF) and flexural strength (rF) values were averaged on 5 samples for each investigated material. Low-velocity impact tests were conducted using an instrumented drop-weight impact testing machine (model Fractovis Plus from CEAST – Italy) equipped with a hemispherical tip (diameter 12.7 mm). All tests were performed putting the sample on a stainless steel annular ring (internal diameter 40 mm, outer diameter 60 mm) and using an impact energy of 31 J (impact parameters: mass = 6.926 kg, velocity = 2.98 m/s). A minimum of 4 samples for each composition, measuring 80 80 mm2 and cut from the prepared laminates, were tested and their mean values and variance were calculated. Optical analysis in reflection mode was carried out by using a microscope (BX51 from Olympus, Japan) to investigate the fibers impregnation in the composite and the distribution of particles (detectable with this optical technique) dispersed in the composite matrices. Fractured surfaces were acquired by a field emission scanning electron microscope, model QUANTA-200FEG from FEI (Eindhoven, The Netherlands). The examined surfaces were coated with a thin layer of a gold-palladium alloy prior to SEM analysis. All glass–fiber composite samples were prepared by polishing the observation surfaces with wet sandpaper and then with a very fine polishing paste.
Table 2 Relevant thermal properties of all matrix compositions. Matrix
Tm,onset (°C)
Tm,peak (°C)
DHm (J/g)
Xc (%)
M–LV M–HV M–LV–PB M–HV–PB
155.8 156.2 154.7 155.7
167.3 169.3 168.1 169.8
69.7 81.9 71.4 78.5
33.7 39.6 34.5 37.9
fiber volume content with respect to the not compatibilized ones (C–LV). On the contrary, C–HV–PB samples exhibited a higher fiber content with respect to the not compatibilized composite (C–HV). This will affect the flexural response of the composite, as shown in the following. Composites were also observed to check the presence of voids within fiber bundles and in the polymer matrix. Impregnation of fabrics resulted to be very good despite of their very different melt flow index, as shown in Fig. 3: no voids were detected by several optical acquisitions in any laminate and all bundles appeared to be well impregnated by the hosting matrix. Tests were performed to measure the percentage of voids with digestion methodology according to the ASTM D3171-99 standard but results were all below the 2% limit of acceptance for the composite (Table 4). Furthermore, the fiber volume content was evaluated according to two different standard (namely ASTM D3171-99 and ASTM D792), evidencing very good accordance between them.
3. Results and discussion 3.2. Flexural properties 3.1. Thermal properties DSC heating thermograms of all matrix compositions are shown in Fig. 2, and relevant thermal properties (melting temperature onset Tm,onset, melting temperature peak Tm,peak, melting enthalpy DHm and relative crystallinity Xc) and summarized in Table 2. All materials showed a single endothermic peak, whose position is only slightly influenced by the inclusion of the compatibilizer, which reduces the melting temperature onset and increases the melting temperature peak. The degree of crystallinity of LV polypropylene was lower than that of HV polypropylene. The addition of compatibilizer marginally affected Xc even if it had an opposite behavior on LV and HV matrices, being increased and decreased, respectively. Density measurements on composites are shown in Table 3. LV compatibilized laminates (C–LV–PB) were characterized by a lower
Fig. 2. DSC heating thermograms of all matrix compositions.
The stress–strain curves from the flexural characterization of matrices are shown in Fig. 4. Neat samples did not break at 5% of strain, and only flexural modulus and strength were evaluated (Table 4). Neat polymers exhibited a higher modulus with respect to compatibilized ones, which acted as a plasticizer. Both flexural modulus and strength of high viscosity polypropylene matrix (M–HV) were the highest among the tested matrices, but its blending with the compatibilizer significantly decreased the bending stiffness. Bending tests performed on composite led to the breaking of samples for the failure of the upper side (glass fiber breakage in compression) as shown in Fig. 5, where broken bundles are evidenced. C–LV samples exhibited lower flexural modulus with respect to the C–HV laminate, because of its lower amount of fiber content and lower matrix stiffness. On the contrary, the flexural strength was significantly higher. The use of compatibilized matrices had positive effect on both flexural modulus and strength (Fig. 6). The improvement is clear in C–LV–PB samples with respect to C–LV due to their matching compositions, having the same fiber volume content. More tricky was the analysis of the C–HV and C– HV–PB data, because the latter was characterized by a slightly higher fiber content. The use of compatibilizer resulted in an improvement of the flexural strength, but it was not clear if the flexural modulus was improved for the better compatibility of fibers with the matrix or for the higher amount of fibers. As also found in the literature [9], the compatibilizer marginally improves the elastic modulus, but it leads to a more pronounced increase of both flexural strength and strain at yield. These results come from the enhanced capability, in compatibilized systems, of load transfer between the matrix and the woven fabric through the interface. The marked damaging of the upper surface of composites after the peak during flexural tests, as exemplified in Fig. 5, comes from the compressive failure of glass fibers, located within the upper layer, occurring at a lower stress in non compatibilized composites with respect to compatibilized ones. After the peak stress, a lowering of the stress-strain curve was detected, which is due to the propagation of both interlaminar and intralaminar damaging of
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Table 3 Density and fiber volume measurements on composites. Matrix
Thickness (mm)
ASTM D 3171-99 3
C–LV C–LV–PB C–HV C–HV–PB
1.30 ± 0.04 1.32 ± 0.07 1.24 ± 0.04 1.20 ± 0.01
ASTM D792
Density (g/cm )
Vf (%)
Void content (%)
Density (g/cm3)
Vf (%)
1.69 ± 0.01 1.69 ± 0.00 1.73 ± 0.02 1.75 ± 0.00
48.99 ± 0.52 48.48 ± 0.32 51.18 ± 1.04 52.39 ± 0.24
1.16 ± 0.17 0.90 ± 0.35 0.46 ± 0.11 0.79 ± 0.49
1.70 ± 0.01 1.69 ± 0.00 1.74 ± 0.02 1.76 ± 0.01
48.46 ± 0.478 47.92 ± 0.150 50.87 ± 1.114 52.17 ± 0.445
Fig. 3. Impregnation of fiber bundles in the different laminate compositions: (A) C–LV, (B) C–LV–PB, (C) C–HV and (D) C–HV–PB.
Table 4 Relevant data evaluated from the flexural tests performed on investigated composites. Sample
Flexural modulus (MPa)
Flexural strength (MPa)
Strain at yield (%)
M–LV M–LV–PB M–HV M–HV–PB C–LV C–LV–PB C–HV C–HV–PB
1149 ± 40 1136 ± 42 1409 ± 92 1329 ± 90 15708 ± 234 16197 ± 405 16276 ± 530 16808 ± 434
26.1 ± 0.5 25.9 ± 0.6 31.7 ± 0.8 30.5 ± 0.8 143.1 ± 3.2 172.9 ± 6.1 129.6 ± 8.2 179.8 ± 9.6
1.43 ± 0.10 1.74 ± 0.09 1.22 ± 0.13 1.62 ± 0.15
Fig. 5. Laminate failure after flexural tests. Fiber bundle breakages are evidenced in the upper side of the sample.
3.3. Low-velocity impact tests
Fig. 4. Stress–strain curves of flexural tests on matrices.
glass fibers and fiber/matrix interfaces, until the failure reaches the lower composite surface.
Low velocity impact tests have been performed by using an impact energy value higher than that bearable by any of the prepared composites, in order to measure the impact strength of the different configurations and to evaluate the force versus deflection and energy versus time curves up to the sample penetration. The impact characterization has shown a different behavior of the structural response with respect to the bending test. In fact, a lower peak force, even if coupled to a higher slope of the force/deflection curve has been exhibited by compatibilized samples with respect to not compatibilized laminates. This has been explained as a result of the increased flexural modulus induced by the presence of compatibilizer. The best performance in terms of dissipated energy has been shown by the C–HV samples, which also reported the highest peak force (Fig. 7). On the contrary, C–HV–PB samples
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the normal forces on the fibers and the friction during fiber slipping justifying the increased energy dissipated until the force peak is reached during impacts on not compatibilized HV samples with respect to not compatibilized LV ones. 3.4. Morphological analysis
Fig. 6. Flexural stress–strain curves for the prepared composites.
exhibited the lowest impact energy absorption after impact, probably as a result of the strong fiber/interface interaction allowed by the compatibilizer. The absorbed energies, evaluated from curves of Fig. 7, are reported as a function of time in Fig. 8. Composites prepared without the compatibilizer absorbed more energy (from 20% to 60%) than their compatibilized ones, as a result of higher load bearing and larger deformations allowed by fibers slipping (Table 5). The strong interaction between fibers and matrix hinders slipping of fibers (fiber pull-out) and it is not able to stop the propagation of the damage. On the contrary, in not compatibilized systems the crack propagation is retarded by the early slipping of the fiber/matrix interface, which results in a post-posed failure of the composite as a whole and in higher deflections and force peaks. Ultimately, the overall laminate toughness of not compatibilized samples resulted to be significantly increased by the fiber pull-out phenomena which is also believed to allow a rearrangement of the fabric to better bear the load, leading to more resilient structures and increased dissipated energies. A slight contribution to the different impact behavior between not compatibilized LV and HV samples could be attributed to their difference in crystallinity. In fact, presumably the higher degree of crystallinity of HV samples with respect to LV ones could lead to higher densification of the surrounding amorphous regions and, consequently, might induce a tightening of the matrix onto the glass fibers. These effects increase
Fig. 7. Force versus deflection curves of composites from the low velocity impact tests.
The analysis performed on the force/deflection curves has been confirmed by SEM micrographs of the fracture surfaces of impacted samples. The improved adhesion of matrix on fibers is evident in Fig. 9B and D, respectively for LV and HV based composites. In laminates produced without the compatibilizer, the fibers exhibit very smooth surfaces, confirming their slippage in the matrix as shown in Fig. 9A and C for LV and HV based composites, respectively. When the high viscosity PP was used, surfaces were very weakly impregnated, as a result of the poorer fiber/matrix interaction. SEM analysis also showed that fracture surfaces of compatibilized composites exhibited a large number of short fibers protruding from the fractured surfaces, unlike not compatibilized ones. This could be an evidence that crack propagation through fibers during impact was very fast and straight through the laminate in the compatibilized composites. On the contrary, in not compatibilized samples the weak interactions between the two phases allowed the pull-out mechanism during impact. The fibers pulling out from the matrix can be considered the main reason for the improved amount of absorbed energy during the impact, because the contribution of the friction (dissipative mechanism allowed by the relative motion of fiber/matrix surfaces) between the sliding fibers and the matrix is not present in composites with compatibilized matrices. This phenomenon resulted more pronounced in the case of high viscosity PP composites (C–HV). It is worth to note that damage zone after impact, depicted in Fig. 10, presents cross shaped cracks in both compatibilized samples, as usually occurs in composites with strong adhesion between fibers and matrix, while in not compatibilized samples damages and cracks are evenly distributed around the impacted hole, not presenting preferential directions of crack propagation along the main directions of the glass fiber fabric. 3.5. Ductility index Further considerations on the impact behavior of the prepared composite systems can be obtained comparing their ductility index (Table 5), defined as the ratio between the energy absorbed after the force peak (Etot Epeak) and the energy absorbed up to the peak
Fig. 8. Absorbed energy versus time curves of composites from the low velocity impact tests.
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Table 5 Data from low velocity impact characterization. Sample
C–LV C–LV–PB C–HV C–HV–PB
Properties at maximum load Load (N)
Deflection (mm)
Epeak
Etot Epeak
Ductility index (DI)
2863 ± 69 2670 ± 52 3228 ± 36 2487 ± 56
5.94 ± 0.14 5.28 ± 0.21 6.81 ± 0.26 5.42 ± 0.26
7.74 ± 0.42 6.36 ± 0.35 9.86 ± 0.38 5.95 ± 0.32
8.71 ± 0.13 8.64 ± 0.71 9.82 ± 1.47 9.53 ± 0.17
1.13 ± 0.07 1.36 ± 0.12 1.00 ± 0.16 1.60 ± 0.06
energy dissipated by mechanisms not involving the breaking of the reinforcing fibers, while after the force peak large fibers breaking occurs. Results demonstrated that the presence of the compatibilizer increases the ductility index by decreasing the area under the force-deflection curve before its maximum with respect to the area under the same curve at high deflections. This behavior seemed to be reversed for not compatibilized systems as a result of the increased energy dissipated before large fiber breaking. In fact, typical dissipative mechanisms, involved in these latter samples and related to the slipping of fibers within the matrix, occur only during the early stage of the impact before fibers failure. In other words, a higher amount of energy is absorbed in the fiber’s elastic region in not compatibilized laminates with respect to compatibilized ones, thanks to the occurrence of pullout phenomena, which in turn seem to occur during the early stage of the impact, before fibers or matrix failures.
4. Conclusions
Fig. 9. Fracture surfaces of impacted samples (A) C–LV, (B) C–LV–PB, (C) C–HV and (D) C–HV–PB.
Fig. 10. Optical pictures of perforated samples (A) C–LV, (B) C–LV–PB, (C) C–HV and (D) C–HV–PB.
force (Epeak), calculated as the integral of the force with respect to the deflection. The first part of the curve represents the amount of
Composite laminates have been produced by using two types of polypropylene matrices, differing for their melt flow index, and a glass fiber fabric. Their static and low velocity impact properties have been evaluated and compared to those exhibited by composite laminates produced by adding a compatibilizing agent to the matrices. The higher amount of crystalline phase, detected by thermal analysis, was responsible for the higher flexural modulus exhibited by the high viscosity PP with respect to the low viscosity one. This consideration also affects laminates behavior and it could have an important contribution on their improved impact behavior. The use of the compatibilizer in both neat matrices and composites resulted in improved static structural properties, namely flexural modulus and flexural strength. While the flexural modulus was only marginally improved by the presence of the compatibilizer (3% increase), the flexural strength was strongly affected, up to an increase of the flexural strength in excess of 30%, in reinforced laminates. The reduced flexural strength in not compatibilized systems has been related to the low capability to transfer the load from the polymeric matrix to the reinforcing fabric. The low velocity impact tests, performed by perforating samples, have shown that the interface between fibers and matrix has a primary role in increasing the impact strength. Force vs deflection curves show that compatibilized systems exhibit higher slopes, as a consequence of their higher stiffness, but incur failure at lower forces with respect to not compatibilized systems. The reduced interface strength in the latter between fibers and matrix (also confirmed by the SEM analysis of composites’ fracture surfaces) was responsible for the occurrence of fibers pull-out phenomena, which dissipate high amount of energy through the friction consequent to sliding mechanisms. The ductility index, evaluated from the impact curves, established that the dissipative mechanisms through relative motion between fibers and matrix were the only reason for the improved impact strength in not compatibilized composites, and that they occurred before fibers failure.
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