- Email: [email protected]
HDPE composites
Materials Science and Engineering A 491 (2008) 492–500
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Isothermal crystallization and mechanical behavior of ionomer treated sisal/HDPE composites Arup Choudhury ∗ Department of Polymer Engineering, Birla Institute of Technology, Mesra, Ranchi 835215, India
a r t i c l e
i n f o
Article history: Received 3 January 2008 Received in revised form 18 February 2008 Accepted 7 March 2008 Keywords: Natural fibre–HDPE composites Isothermal crystallization Surface morphology Mechanical properties Analytical modelling
a b s t r a c t Isothermal crystallization behavior and mechanical properties of ionomer-treated sisal fibres reinforced high-density polyethylene (HDPE) composites have been investigated. Avrami analysis of the crystallization kinetics showed high crystallization rate (Kn ) and short crystallization half time (t0.5 ) for composites compared to neat HDPE, indicating strong nucleating ability of sisal fibres. The crystallization activation energy (Ea ) of the composites decreased with increasing fibre content. The polarizing optical microscopy showed the occurrence of transcrystallization at the fibre–matrix interface. The HDPE/sisal composites displayed superior tensile and flexural properties compared to neat HDPE. For these composites, predictions of Young’s modulus using the Modified Rule of Mixtures and Cox model correlated well with the experimental data. © 2008 Elsevier B.V. All rights reserved.
1. Introduction In recent years, natural fibres have found use as potential resources for making high-performance composite material. The attention in natural fibre reinforced polymer composites is growing rapidly due to their high mechanical performance, wide versatility, good processing advantages, low cost, and low density [1,2]. Lignocellulosic fibres such as jute, sisal, hemp, coir, and banana have been successfully used as reinforcing materials in many polymeric resin matrices [3–6]. Among these fibres, sisal (Agave sisalana) is of particular interest in that its composites have high impact strength besides having moderate tensile and flexural properties compared to other lignocellulosic fibres [7]. However, some studies have been reported in the literature on the use of sisal fibre as a reinforcing agent in thermoplastic matrices [8–13]. The major drawback of these natural fibre/thermoplastic composites is the inherent incompatibility of the hydrophilic lignocellulosic fibres with the hydrophobic thermoplastic matrices, which calls for improving the fibre–matrix interfacial adhesion by using suitable compatibilisers and coupling agents [14–16]. Among many coupling agents currently used for improvement of the fibre–matrix interfacial bond strength, ionomers based on copolymers of acids and olefin monomer units are worthy of exploration. The amphiphilic nature may allow these ionomers to be
∗ Tel.: +91 9430 732461; fax: +91 651 2276184. E-mail address: [email protected]. 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.03.011
compatible with both the matrix and fibre, and thus could act as coupling agents [17,18]. For composites based on semicrystalline polymers, the morphological features like degree of crystallinity, spherulite size, lamella thickness and crystallite orientation play an important role in influencing the ultimate properties of the polymer matrix, and thus the composite. It was found that both morphology and crystallinity of thermoplastic polymer matrices deeply affected by fibre reinforcement [19–21]. In the case of thermoplastic composites, the growth of highly oriented transcrystalline layers at fibre–matrix interphase is thought to have a critical influence on improvement of the mechanical properties, due to better fibre–matrix adhesion [22]. The occurrence of transcrystallization depends on the type of fibre used and the crystallization temperature [23,24]. A transcrystalline layer can be formed at the fibre–matrix interface regions if the fibres with high nucleating ability were employed. However, there were not many studies devoted to investigate the crystallization behavior and microstructure of the natural fibre reinforced thermoplastic composites [23–28]. The mechanical properties of fibre-reinforced composites strongly depend on several factors such as fibre size, fibre loading, fibre dispersion, fibre orientation and fibre–matrix interfacial bond strength [29,30]. In many occasions, the effects of fibre reinforcement on the mechanical properties have been explained by comparing the experimental results with existing micromechanical models such as Parallel and Series, Hiesch, Cox, Halpin–Tasi, Modified Halpin–Tasi, Bowyer–Bader models, etc. [31–35]. It was
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
493
Table 1 Physical, chemical and mechanical properties of sisal fibres
Table 2 Formulation of the different HDPE/sisal composites
Fibre
Composites
Average diameter (m) Average length (mm) Density (g/cm3 ) Cellulose (%) Hemi-cellulose (%) Lignin (%) Tensile strength (MPa) Tensile modulus (MPa) Elongation (%)
Sisal 100 4.0 1.3–1.5 56.52 16.49 10.62 580 773 4.3
(95:5) HDPE/fibre (90:10) HDPE/fibre (85:15) HDPE/fibre (80:20) HDPE/fibre
Weight percentage (wt%) HDPE matrix
Ionomer-treated sisal fibre
Volume fraction of fibre, Vf (%)
95 90 85 80
5 10 15 20
3.22 6.44 9.66 12.88
2.4. Composites fabrication
concluded that the mechanical properties of fibre-reinforced composites strongly depend upon the fibre size distribution and fibre orientation distribution. However, a limited number of research works was dedicated to assess the viability of experimental data by using theoretical models for randomly oriented short fibre reinforced composites. The aim of this work is to investigate the effect of sisal fibre reinforcement on the crystallization kinetics and morphology of high-density polyethylene (HDPE). Surlyn ionomer was used as a coupling agent to improve the composites performance. The variations of tensile, flexural and impact properties as a function of fibre contents were studied. Moreover, a comparison between experimental and estimated Young’s modulus using various micromechanical models was presented for the HDPE/sisal composite.
2. Experimental 2.1. Materials The pure high-density polyethylene (M6805U: melt flow index = 0.5 g/10 min at 190 ◦ C and density = 0.968 g/cm3 at 23 ◦ C) was obtained from Haldia Petrochemical Ltd. India. The coupling agent used in this investigation was Surlyn® 1650 (density = 0.94 g/cm3 ; melt flow index = 1.5 g/10 min at 190 ◦ C; melting temperature = 97 ◦ C with a heating rate of 10 ◦ C/min; methacrylic acid co-monomer content = 3.7–4.0 mol% and 57 mol% acid was neutralized by sodium salt), procured from Du Pont USA. Sisal fibres (A. sisalana) were collected form local sources. Table 1 presents the physical, chemical and mechanical properties of sisal fibres. Sodium hydroxide (NaOH) and acetic acid used for chemical treatment of the fibres were obtained from E. Merck, India.
2.2. Washing and alkali treatment
Composites were prepared by the melt mixing of the ionomertreated short sisal fibres with HDPE in a HAAKE rotor mixture (Model Rheomix 600, Dreieich, Germany) at 120 ◦ C with roller blades and a mixing chamber with a 60 cm3 volumetric capacity. The process was carried out for 10 min at an optimum speed of 50–60 rpm, which depended on the quantity of the fibres by weight. Each batch contained various plastic to fibre weight ratio (95:5, 90:10, 85:15 and 80:20). The composite formulations are shown in Table 2. Each composite mixture was then homogenized in a tworoll mill (150E-400 Collins, Germany) at 130 ◦ C and compression molded with a Delta Malikson 100TY pressman (Mumbai, India) to produce composite sheets (3 ± 0.2-mm thick). Test specimens were prepared from these sheets according to ASTM (American Society for Testing and Materials) standard D with a counter-cut copy-milling machine (6490, Ceast, Italy) with calibrated templates. 2.5. Thermal analysis The thermal behavior of the neat HDPE and formulated HDPE/sisal composites was measured with a differential scanning calorimetric (DSC) analyzer (DSC-10Q, TA Instruments, New Castle, DE, USA) under nitrogen atmosphere. The samples were heated to 170 ◦ C with a heating rate of 10 ◦ C/min and kept at this temperature for 10 min to eliminate previous thermal history and then rapidly cooled (at 100 ◦ C/min) to crystallization temperature (Tc ), and maintained at Tc till the time necessary for complete crystallization of the polymer matrix. The heat evolved during the isothermal crystallization (Hc ) was recorded as a function of time (t), at different crystallization temperatures. After crystallization, the samples were heated to melting point at a rate of 10 ◦ C/min. The crystalline melting temperatures (Tm ) and degree of crystallinity (Xcr ) of the samples were determined from the endotherm. The enthalpy of crystallization of 100% crystalline polyethylene was taken as 290 J/g [36].
The short sisal fibres (length ∼4.0 mm) were thoroughly washed by aqueous detergent solution to remove dirt followed by cleaning with distilled water and dried in vacuum oven at 70 ◦ C for 1 day. The washed fibres were immersed in 5% aqueous NaOH solution for a period of 1 h at 40 ◦ C, and this was followed by washing with 0.1N acetic acid and distilled water. The alkali-treated fibres were than dried in vacuum oven at 70 ◦ C for 2 days to obtain mercerized fibres.
2.6. Mechanical properties
2.3. Coupling agent treatment
2.7. Morphological analysis
The mercerized sisal fibres were immersed in hot 5% Surlyn ionomer solution (in toluene) at 70 ◦ C for 30 min to obtain ionomertreated fibres.
The surface morphology of the composites was analyzed using a polarized optical microscopy (Leika Metalographic Aristomet model), which attached with automatic hot-stage thermal
Rectangular bar specimens for tensile and flexural tests were prepared according to ASTM D638 and ASTM D790, respectively. The measurements of the samples were conducted with an Instron 3366 tensile test machine (Grove City, PA, USA) at a crosshead speed of 2 mm/min. The Izod impact strength of notched samples (2 mm depth) was measured with a Davenport Izod impact tester as per ASTM D256 method.
494
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
Table 3 Melting and crystallization kinetic parameters of neat HDPE and HDPE/sisal composites obtained from DSC measurements Material HDPE/fibre
Thermal properties Tc a (◦ C)
nb
Kn c
t0.5 d (s)
Tm e (◦ C)
Hm f (J/g)
Eg (kJ/mole)
% Xcr h
100/0 100/0 100/0 100/0
116 119 121 124
2.06 2.19 2.43 2.62
0.05643 0.00255 0.00021 0.00003
374 868 1288 1405
126.6 127.1 129.8 131.4
82.94 86.42 89.03 90.48
85.8
28.6 29.8 30.7 31.2
95/5 95/5 95/5 95/5
116 119 121 124
2.32 2.51 2.74 2.93
0.14148 0.01014 0.00091 0.00008
254 628 726 945
124.7 125.5 127.6 130.4
86.42 90.48 91.35 96.86
108.3
29.8 31.2 30.5 33.4
90/10 90/10 90/10 90/10
116 119 121 124
2.28 2.79 2.83 2.88
0.46613 0.01588 0.00179 0.00017
232 511 610 855
123.7 124.4 126.4 128.7
92.22 95.41 97.15 100.63
127.6
31.8 33.9 34.5 34.7
85/15 85/15 85/15 85/15
116 119 121 124
2.68 2.73 2.96 3.01
0.52967 0.02548 0.00231 0.00016
208 458 485 642
123.3 123.8 125.5 127.9
98.6 102.66 107.01 109.04
134.8
33.0 33.4 36.9 37.6
80/20 80/20 80/20 80/20
116 119 121 124
2.85 2.93 3.25 3.51
0.72664 0.02734 0.00355 0.00033
153 376 422 541
122.8 124.5 125.3 127.6
108.87 110.35 117.79 116.49
151.5
36.5 37.0 39.5 39.0
a b c d e f g h
Crystallization temperature. Avrami exponent. Crystallization kinetic constant. Crystallization half time. Crystalline melting temperature. Heat of fusion. Crystallization activation energy. Crystallinity.
unit (Mettler FP-90). The composite samples were sandwiched between microscope glass slides and inserted into the hot stage unit. The samples were heated to 170 ◦ C at 10 ◦ C/min, held at this temperature for 10 min, then rapidly cooled to 116 ◦ C at 100 ◦ C/min and allowed to crystallize isothermally. Images were obtained using a Nikon digital camera connected to the microscopy.
2.8. Fourier transform infrared (FT-IR) spectroscopy Infrared spectroscopic analysis of the ionomer-treated and untreated sisal fibres was carried out using an FT-IR Thermo Nicolet (NEXUS 870) spectrophotometer at 25 ◦ C.
3. Results and discussion 3.1. Melting and crystallization behavior of HDPE/sisal composites The quantitative results of melting and crystallization behavior as determined from DSC measurements are summarized in Table 3. The crystalline melting temperature of HDPE matrix in the composites was recorded lower compared to neat HDPE. The Tm values in the composites decreased with increasing fibre content (Table 3). The occurrence of lower melting peak in the composites could be attributed to the fact that strong nucleation on fibre surfaces shortened the time required for HDPE crystallization, therefore, limiting the isothermal thickening of HDPE crystals and decreasing their melting temperature. For all examined samples, an increase
Fig. 1. Heat flow curves of (a) neat HDPE and (b) 80:20 HDPE/sisal composite during isothermal crystallization.
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
495
Fig. 2. Plots of relative crystallinity (Xt ) versus time (t) at various Tc obtained for (a) neat HDPE, (b) 90:10 HDPE/sisal composites and (c) 80:20 HDPE/sisal composites.
of the melting temperature (Tm ) with increasing Tc was observed. The overall crystallization rate (1/t0.5 ) at higher Tc was found to slow (Table 3), which might be caused the formation of thick HDPE crystals having higher melting temperature. The percent crystallinity of HDPE (%Xcr ) was significantly increased by fibre reinforcement, indicating strong heterogeneous nucleating ability of the sisal fibres. In addition, the nucleating effects of the fibre surfaces promoted the growth and formation of interfacial transcrystalline layer when the composite sample was cooled from the melt, as observed in the optical microscopy images [37,38]. The effect of sisal fibre reinforcement on the crystallization behavior and thermodynamics of HDPE matrix was evaluated by analyzing the isothermal crystallization kinetics of neat HDPE and its composites in the temperature range of 116–124 ◦ C. Fig. 1 shows the heat flow curves of 100% HDPE and 20% fibre filled HDPE/sisal composite during isothermal crystallization at various Tc . It was seen from Fig. 1 that at any crystallization temperature, the primary crystallization of HDPE matrix in the composites was almost completed before the crystallization of neat HDPE had reached to 50–70% conversion. The fast crystallization conversion in the composites could be attributed to the transcrystallization as a result of nucleating effects of fibre surfaces. The short induction time and high crystallization rate are supporting this hypothesis. The plots of relative crystallinity (Xt ) as function of time t at various crystallization temperatures (Tc = 116, 119, 121 and 124 ◦ C) are shown in Fig. 2. From the comparison among the crystallization isotherms at the same crystallization temperature it can be clearly pointed out that the crystallization rate of HDPE matrix in the composites was faster than that of neat HDPE, indicating the occurrence of
nucleating effects of sisal fibres in the composites. Fig. 3 illustrates the dependence of overall crystallization rate (Rc ), expressed as the reciprocal of crystallization half time (1/t0.5 ), on the crystallization temperature for all examined samples. The crystallization rate (Rc ) of composites was recorded higher compared to neat polymer at the same Tc , and it can be observed to increase with increasing fibre content (Fig. 3). One possible reason was that strong nucleation on
Fig. 3. Reciprocal of crystallization half time (t0.5 ) versus Tc for neat HDPE and different HDPE/sisal composites.
496
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
Fig. 4. Avrami plots of the (a) neat HDPE, (b) 90:10 HDPE/sisal composites and (c) 80:20 HDPE/sisal composites at various Tc .
fibre surfaces shortened the time required for HDPE crystallization. For all examined samples, the crystallization process was hindered with increasing Tc , as anticipated on the basis of the nucleation theory [39]. 3.2. Isothermal crystallization kinetics The crystallization kinetics of the neat polymer and its sisal fibre composites are analyzed by using the Avrami model [40]: Xt = 1 − exp(−Kn t n )
(1)
The linearisation of Eq. (1) leads to Eq. (2): ln[− ln(1 − Xt )] = n ln t + ln Kn
Avrami exponent in the composites was varied around 3.0 implying a heterogeneous nucleation with three-dimensional growth of HDPE crystals. In contrast, the value of n was varied about 2.0 for the neat HDPE indicating two-dimensional diffusion controlled crystal growth. At any crystallization temperature, the crystallization rate constant of HDPE exhibited a significant increment as the fibre content increase (Table 3). The higher crystallization rate constant of the composites could be ascribed to the nucleating effect of the fibres. The crystallization rate constant Kn can be presented by Arrhenius equation: 1/n
Kn (2)
where Kn and n are the crystallization rate constant and Avrami exponent, respectively. The Avrami approach is the most applied tool to describe the phenomenology of the crystal growth. The Avrami exponent depends on growth geometry and nucleation type of the crystals. According to Eq. (2), the values of Kn and n were obtained from the intercept and slope of the linear plots of ln[−ln(1 − Xt )] versus ln(t). Fig. 4 demonstrates the Avrami plots of the neat HDPE, (90:10) HDPE/sisal and (80:20) HDPE/sisal composites at various crystallization temperatures. For all examined samples, the values of Avrami parameters (n and Kn ) in the temperature range 116–124 ◦ C are complied in Table 3. As observed in Table 3 that n increased with increasing Tc , and it recorded slightly higher for the HDPE/sisal composites compared to neat HDPE. The
= K0 exp
−E RTc
(3)
where K0 , R and E are the pre-exponential factor, universal gas constant and activation energy, respectively. The value of activation energy for the crystallization process was calculated from the slopes of the linear plots of (1/n) ln Kn versus 1/Tc . The values of E of neat HDPE and composites are summarized in Table 3. The activation energy of HDPE crystallization was found to be significantly enhanced by the addition of short sisal fibre to HDPE matrix. The crystallization activation energy is the sum of the transport activation energy and the nucleation activation energy. As mentioned above, the nucleation effect of sisal fibres leads to enhance the crystallization of HDPE in the composites. Hence, the higher crystallization activation energy of the composite might be originated from their greater transport activation energy because the mobility of HDPE chain segments was strongly restricted
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
497
Fig. 5. Polarized optical micrographs of the isothermally crystallized 80:20 HDPE/sisal composites (Tc = 116 ◦ C).
Table 4 Mechanical properties of HDPE/sisal composites at different fibre content Sample
Mechanical properties Tensile strength (MPa)
Pure HDPE (95:5) HDPE/fibre (90:10) HDPE/fibre (85:15) HDPE/fibre (80:20) HDPE/fibre
26.0 28.5 30.7 32.0 32.8
± ± ± ± ±
0.5 1.9 1.8 2.3 2.8
Tensile modulus (MPa) 860 952 1086 1107 1219
± ± ± ± ±
0.3 2.3 3.7 3.3 3.4
by the fibres and consequently hindered the crystallization process. 3.3. Optical microscopy study As observed in the polarized optical microscopic analysis, the growth and morphology of the HDPE crystals were significantly influenced by the addition of sisal fibres. Fig. 5 shows the optical photomicrographs of the isothermally crystallized (80:20) HDPE/sisal composites (Tc = 116 ◦ C). As shown in the photomicrographs, the nucleation of HDPE spherulites as well as regular transcrystalline growth occurred along the fibre–matrix interfaces. However, the formation and growth of transcrystallinity in a composite depend upon the performances of fibre–matrix interfaces [41]. As seen in Fig. 5, the crystalline morphology at the fibre–matrix interface does not look like the morphology of the crystals in the bulk of the matrix. In this situation, the transcrystalline layer was grown isothermally at the fibre–matrix interface while the bulk of the matrix still remained in the molten form, thus the growth transcrystallinity was practically undisturbed. The thickness of the transcrystalline layer was roughly constant of about 30 m. The fibre–matrix interface region was made of needle-like compact crystals propagated vertically on the surface of fibre. In earlier investigations, the needle-like crystals have also been identified on polyethylene with a certain molecular weight range under isothermal conditions [42,43]. 3.4. Mechanical properties The properties of matrix and fibre are very important aspects in achieving good mechanical properties of the composites. The melt strength and processability are the most sensitive parameters to matrix polymer, whereas the modulus is dependent on fibre properties [44]. The mechanical properties (tensile strength, tensile modulus, flexural strength, flexural modulus, and impact strength) of neat HDPE and its formulated sisal fibre composites are summarized in Table 4. The composites have shown scattered response to mechanical properties, as presented in Table 4, because of diffi-
Flexural strength (MPa) 37.0 40.7 43.2 45.0 46.7
± ± ± ± ±
0.8 1.5 1.2 1.9 1.9
Flexural modulus (MPa) 928 1131 1194 1354 1468
± ± ± ± ±
0.3 2.7 3.7 2.9 3.5
Impact strength (J/m) 37.0 20.3 17.7 15.2 14.3
± ± ± ± ±
0.16 0.19 0.24 0.22 0.26
culty in proper dispersion of short sisal fibres and random fibre alignment in the matrix. However, all the composites exhibited improved mechanical properties compared to neat HDPE. When the fibre content raised up to 20%, the tensile strength, tensile modulus, flexural strength and flexural modulus of the HDPE/sisal composite increased about 27, 41, 26, and 58%, respectively. The strong interfacial adhesion between HDPE and ionomer-treated fibres lead to improve the stress transfer from the matrix to the fibres, which caused better mechanical performance. At lower fibre content (<10%) the dispersion of fibres was very poor that resulted poor stress transfer, whereas a strong tendency of the fibre–fibre interaction at higher fibre content (>15%) significantly enhanced the stress transfer. A hypothetical model of the interface between ionomer-treated sisal fibre and HDPE is shown in Scheme 1. The PE chain of ionomer diffuses into the HDPE matrix through interchain entanglements. On the other hand, the acrylic groups of ionomer molecules form ester bonds with the hydroxyl groups of the fibres. FT-IR spectra of ionomer-treated and untreated sisal fibres are reported in Fig. 6. The spectra of ionomer-treated sisal fibres show a significant absorption peak in the carbonyl region (1715 cm−1 ), which is associated with the stretching vibration of methacrylic ester bond.
3.5. Modelling of mechanical property The elastic properties of the HDPE/sisal composites can be theoretically determined or derived by using a variety of micromechanical composite models. These models are derived based on the properties of the individual components of the composite and their arrangement [45]. Thus, a micromechanical model might be used to find the best combination of constituent materials to satisfy material design considerations. In this study, four different micromechanical composite models have been compared with experimental results in order to assess the influences of fibre dispersion on the Young’s modulus of HDPE/sisal composites. The micromechanical models of the Young’s modulus for randomly distributed short fibre-based polymer composites are as follows.
498
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
Scheme 1. Hypothetical model of fibre modification with Surlyn ionomer and reaction with HDPE matrix.
3.5.1. Voigt–Reuss model The Voigt–Reuss model [46,47] has been used to estimate the Young’s modulus for a lamina of short random fibre composites. This model is independent on the geometry of reinforcement. It can be expressed as: εVR =
3 5 [V ε + (1 − Vf )εm ] + 8 f f 8
εf εm εf (1 − Vf ) + εm Vf
(4)
where εf and εm are the Young’s moduli of fibre and matrix, respectively. Vf is the volume fraction of fibre in the composites, calculated using the following equation: Vf =
Wf m Wf f + (1 − Wf )f
3.5.2. Halpin–Tsai model The semi-empirical equation developed by Halpin and Tsai [48] has been widely used for predicting the elastic properties of a lamina of short random fibre composites. The following form of the Halpin and Tsai equation was used to determine the Young’s modulus of the composites: εHT =
8
1 − L Vf
+
5 1 + 2T Vf εm 8 1 − T Vf
(6)
where lf and df are the average length and diameter of the fibres. In Eq. (6) the parameters L and T are given as:
(5)
where Wf , f and m denote the weight fraction of fibres in the composite, the density of fibre and the density of matrix, respectively.
3 1 + 2(l /d ) (1 − V ) f f L f
L =
(εf /εm ) − 1 (εf /εm ) + 2(lf /df )
(7)
T =
(εf /εm ) − 1 (εf /εm ) + 2
(8)
3.5.3. Modified Rule of Mixtures The Rule of Mixtures can be used to determine the Young’s modulus for three-dimensional random short fibre composites [49,50]. This model works extremely well for aligned continuous fibre composites. The Young’s modulus was determined by using this equation: εMRM =
1 V ε + (1 − Vf )εm 5 f f
(9)
3.5.4. Cox model This model introduces the influence of the reinforcement aspect ratio into the expression of Rule of Mixtures [51]. The Young’s modulus of three-dimensional random short fibre composites was evaluated using this equation:
Fig. 6. FT-IR spectra of ionomer-treated and untreated sisal fibres.
εCox =
1 L εf Vf + (1 − Vf )εm 5
(10)
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
499
Table 5 Experimental and predicted Young’s modulus for neat HDPE, sisal fibre and HDPE/sisal composites Sample
Young’s modulus, ε (GPa) Experimental
Pure HDPE Sisal fibre (95:5) HDPE/fibre (90:10) HDPE/fibre (85:15) HDPE/fibre (80:20) HDPE/fibre
1.195 12.40 2.002 2.425 2.975 3.124
± ± ± ± ± ±
0.1 2.7 0.2 0.3 0.2 0.3
Voigt–Reuss model
Halpin–Tsai model
Modified mixture law
Cox model
– – 3.333 4.688 5.439 5.947
– – 3.003 4.115 4.652 5.277
– – 2.313 3.033 3.553 3.771
– – 2.207 2.544 3.037 3.281
indicating that the assumptions made to derive these theoretical equations were closely agreeing with experimental conditions. The relative deviation between experimental and theoretical results increased with increasing fibre content, but it was well within agreeable limits for the Modified Rule of Mixtures and the Cox model. This suggested that the Modified Rule of Mixtures and Cox model can be applied to predict Young’s modulus of HDPE/sisal composites. However, the best theoretical–experimental correlation was obtained with the Cox relationship because it takes into account the fibre aspect ratio. Apart from that, the modulus of composite depends on the homogeneity of fibre dispersion at a given volume fraction of sisal fibre. Obviously, the higher modulus was obtained for better fibre distribution. The composite Young’s modulus was nearly the matrix modulus for poor fibre dispersion. 4. Conclusions
Fig. 7. Comparison of experimental Young’s modulus values of HDPE/sisal composites with different micromechanical models.
The independent parameters of Eq. (10) are calculated using the following equations:
L = 1 − s=
tan(ˇs) ˇs
2lf 0.5df
ˇ2 =
2εm εf (1 + m ) ln(1/Vf )
(11)
(12) (13)
where m is the Poisson’s ratio of HDPE matrix. Generally, Young’s modulus of the composite material reflects the capability of both fibre and matrix to transfer the elastic deformation under small strain without interface fracture. Table 5 presents the predicted Young’s modulus values obtained from micromechanical models and the experimental data for all formulated samples. It can be seen that the Young’s modulus of the composites increased as the fibre volume fraction increase. A comparison between the experimental data and predicted Young’s modulus values is demonstrated in Fig. 7. The predicted Young’s modulus values were recorded higher than experimental values (Fig. 7). Fig. 7 shows a greater degree of divergence between the experimental Young’s modulus values and the predicted values using Voigt–Reuss and Halpin–Tsai models, because these models were developed for lamina composite (two-dimensional structure), not for three-dimensional composite. The differences of these two micromechanical models with the experimental Young’s modulus were recorded ∼65% for the Voigt–Reuss model and ∼50% for the Halpin–Tsai model. Although, the comparison between experimental Young’s modulus and those obtained from Modified Rule of Mixtures or Cox model demonstrated a close agreement within 10%,
The present investigation showed the usefulness of sisal fibre as a good reinforcing agent and Surlyn ionomer as an effective coupling agent for composite fabrication with polyethylene matrix. DSC analysis demonstrated that the presence and the concentration of short sisal fibres affect markedly the crystallization behavior of HDPE matrix. A noticeable decrease in the half time of HDPE crystallization was observed by the incorporation of short sisal fibre to HDPE. This could be attributed to the fact that the surfaces of sisal fibres act as nucleating sites for the crystallization of the matrix polymer, promoting the growth and formation of transcrystalline layer in the composites. The concentration of nucleating sites in the HDPE/sisal composites increased as the fibre content increase, which was reflected in the increase of the crystallinity of the composites. The Avrami exponent in the composite samples was varied around 3.0 implying three-dimensional growth of HDPE crystals, whereas the n-values of about 2.0 for the neat HDPE suggest two-dimensional diffusion controlled crystal growth. A hot-stage polarized optical microscopy was used to follow the development of transcrystalline layer along the fibre–matrix interface. The mechanical properties of HDPE/sisal composites were found to increase with increasing fibre content. This was ascribed to the reinforcing effect imparted by the ionomer-treated fibres, which allowed a uniform stress distribution from the polymer matrix to the dispersed fibre phase. A comparison between the experimental results and the theoretically predicted Young’s modulus has been presented. Modified Rule of Mixtures and Cox model give more close agreement to experimental data for HDPE/sisal composite. References [1] K.G. Satyanarayana, B.C. Pai, K. Sukumaran, S.G.K. Pillai, Handbook of Ceramics and Composites, Marcel Decker, New York, 1990, pp. 339–360. [2] K.G. Satyanarayana, K. Sukumaran, P.S. Mukherjee, C. Pavithran, S.G.K. Pillai, Natural Fibre-Polymer Composites, Haworth Press, England, 1990, pp. 117–136. ´ adi, ´ ´ G. Nagy, J. Mocz ´ o, ´ B. Pukanszky, ´ [3] L. Dany K. Renner, Z. Szabo, Polym. Adv. Technol. 17 (2006) 967–974. [4] S. Mohanty, S.K. Verma, S.K. Nayak, Compos. Sci. Technol. 66 (2006) 538–547.
500
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
[5] I. Baroulaki, J.A. Mergos, G. Pappa, P.A. Tarantili, D. Economides, K. Magoulas, C.T. Dervos, Polym. Adv. Technol. 17 (2006) 954–966. [6] T. Li, N. Yan, Compos. A 38 (2007) 1–12. [7] C. Pavithran, P.S. Mukharjee, M. Brahmakumar, A.D. Damodaran, J. Mater. Sci. Lett. 6 (1987) 882–884. [8] K. Jayaraman, Compos. Sci. Technol. 63 (2003) 367–374. [9] X. Colom, F. Carrasco, P. Pages, J. Canavate, Compos. Sci. Technol. 63 (2003) 161–169. [10] P.A. Sreekumar, K. Joseph, G. Unnikrishnan, S. Thomas, Compos. Sci. Technol. 67 (2007) 453–461. [11] S. Mohanty, S.K. Nayak, Mater. Sci. Engg. A443 (2007) 202–208. [12] N. Chand, D. Jain, Compos. A 36 (2005) 594–602. [13] M. Idicula, A. Boudenne, L. Umadevi, L. Ibos, Y. Candau, S. Thomas, Compos. Sci. Technol. 66 (2006) 2719–2725. [14] M. Brahmakumar, C. Pavithran, R.M. Pillai, Compos. Sci. Technol. 65 (2005) 563–569. [15] J.M. Park, S.T. Quang, B.S. Hwang, D.V.K. Lawrence, Compos. Sci. Technol. 66 (2006) 2686–2699. [16] L.M. Arzondo, C.J. Perez, J.M. Carella, Polym. Engg. Sci. 45 (2005) 613–621. [17] H. Akimoto, T. Kanazawa, M. Yamada, S. Matsuda, G.O. Shonaike, A. Murakami, J. Appl. Polym. Sci. 81 (2001) 1712–1720. [18] E.R. Sadiku, R.D. Sanderson, Macromol. Mater. Engg. 286 (2001) 472–479. [19] M. Arroyo, R. Zitzumbo, F. Avalos, Polymer 41 (2000) 6351–6359. [20] E. Piorkowska, Macromol. Symp. 169 (2001) 143–148. [21] H. Quan, Z.-M. Li, M.-B. Yang, R. Huang, Compos. Sci. Technol. 65 (2005) 999–1021. [22] H. Schonhorn, F.W. Ryan, J. Polym. Sci. 6 (1968) 231–240. [23] C. Wang, L.M. Hwang, J. Polym. Sci. B 34 (1996) 1435–1442. [24] C. Wang, C.R. Liu, Polymer 38 (1997) 4715–4718. [25] P.V. Joseph, K. Joseph, S. Thomas, C.K.S. Pillai, V.S. Prasad, G. Groeninckx, M. Sarkissova, Compos. A 34 (2003) 253–266. [26] P.V. Joseph, K. Joseph, S. Thomas, Compos. Sci. Technol. 59 (1999) 1625–1640. [27] J.L. Thomason, A.A. Van Rooyen, J. Mater. Sci. 27 (1992) 889–896.
[28] M. Pracella, D. Chionna, I. Anguillesi, Z. Kulinski, E. Piorkowska, Compos. Sci. Technol. 66 (2006) 2218–2230. [29] A.K. Rana, A. Mandal, S. Bandyopadhyay, Compos. Sci. Technol. 63 (2003) 801–806. [30] A.P. Kumar, R.P. Singh, B.D. Sarwade, Mater. Chem. Phys. 92 (2005) 458–469. [31] A.G. Facca, M.T. Kortschot, N. Yan, Compos. Sci. Technol. 67 (2007) 2454– 2466. [32] R.G. de Villoria, A. Miravete, Acta Mater. 55 (2007) 3025–3031. ¨ [33] T.T.L. Doan, S.L. Gao, E. Mader, Compos. Sci. Technol. 66 (2006) 952–963. [34] G. Kalaprasad, K. Joseph, S. Thomas, J. Mater. Sci. 32 (1997) 4261–4267. [35] G. Kalaprasad, K. Joseph, S. Thomas, J. Compos. Mater. 31 (1997) 509–527. [36] J. Brandup, E.H. Immergut, Polymer Handbook, Wiley, New York, 1988 (chapter 8). [37] Y. Huang, J. Petermann, Polym. Bull. 36 (1996) 517–524. [38] A. Amash, P. Zugenmaier, Polymer 41 (2000) 1589–1596. [39] J.D. Hoffman, G.T. Davis, J.I. Lauritzen, The Rate of Crystallization of Linear Polymers with Chain Folding, Crystalline and Noncrystalline Solids, Plenum Press, New York, 1976, p. 497. [40] M. Avrami, J. Chem. Phys. 7 (1939) 1103–1112. [41] S. Incardona, C. Migliaresi, D. Wagner, A. Gilbert, G. Marom, Compos. Sci. Technol. 47 (1993) 43–50. [42] L. Mandelkern, M. Glotin, R.A. Benson, Macromolecules 14 (1981) 22–34. [43] J. Maxfield, L. Mandelkern, Macromolecules 10 (1977) 1141–1153. [44] D. Hull, An Introduction to Composite Materials, 2nd ed., Cambridge University Press, Cambridge, 1996, pp. 326–345. [45] R.M. Jones, Mechanics of Composite Material, Hemisphere Publishing Corporation, New York, 1975, pp. 86–104. [46] W. Voigt, Ann. Phys. 38 (1889) 573–587. [47] A. Reuss, Z. Angrew. Math. Mech. 9 (1929) 49–58. [48] J.C. Halpin, J. Compos. Mater. 3 (1969) 732–734. [49] H. Fukuda, T.W. Chou, J. Mater. Sci. 16 (1981) 1088–1096. [50] H. Fukuda, T.W. Chou, J. Mater. Sci. 17 (1982) 1003–1011. [51] H.L. Cox, Brit. J. Appl. Phys. 3 (1952) 72–79.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Isothermal crystallization and mechanical behavior of ionomer treated sisal/HDPE composites Arup Choudhury ∗ Department of Polymer Engineering, Birla Institute of Technology, Mesra, Ranchi 835215, India
a r t i c l e
i n f o
Article history: Received 3 January 2008 Received in revised form 18 February 2008 Accepted 7 March 2008 Keywords: Natural fibre–HDPE composites Isothermal crystallization Surface morphology Mechanical properties Analytical modelling
a b s t r a c t Isothermal crystallization behavior and mechanical properties of ionomer-treated sisal fibres reinforced high-density polyethylene (HDPE) composites have been investigated. Avrami analysis of the crystallization kinetics showed high crystallization rate (Kn ) and short crystallization half time (t0.5 ) for composites compared to neat HDPE, indicating strong nucleating ability of sisal fibres. The crystallization activation energy (Ea ) of the composites decreased with increasing fibre content. The polarizing optical microscopy showed the occurrence of transcrystallization at the fibre–matrix interface. The HDPE/sisal composites displayed superior tensile and flexural properties compared to neat HDPE. For these composites, predictions of Young’s modulus using the Modified Rule of Mixtures and Cox model correlated well with the experimental data. © 2008 Elsevier B.V. All rights reserved.
1. Introduction In recent years, natural fibres have found use as potential resources for making high-performance composite material. The attention in natural fibre reinforced polymer composites is growing rapidly due to their high mechanical performance, wide versatility, good processing advantages, low cost, and low density [1,2]. Lignocellulosic fibres such as jute, sisal, hemp, coir, and banana have been successfully used as reinforcing materials in many polymeric resin matrices [3–6]. Among these fibres, sisal (Agave sisalana) is of particular interest in that its composites have high impact strength besides having moderate tensile and flexural properties compared to other lignocellulosic fibres [7]. However, some studies have been reported in the literature on the use of sisal fibre as a reinforcing agent in thermoplastic matrices [8–13]. The major drawback of these natural fibre/thermoplastic composites is the inherent incompatibility of the hydrophilic lignocellulosic fibres with the hydrophobic thermoplastic matrices, which calls for improving the fibre–matrix interfacial adhesion by using suitable compatibilisers and coupling agents [14–16]. Among many coupling agents currently used for improvement of the fibre–matrix interfacial bond strength, ionomers based on copolymers of acids and olefin monomer units are worthy of exploration. The amphiphilic nature may allow these ionomers to be
∗ Tel.: +91 9430 732461; fax: +91 651 2276184. E-mail address: [email protected]. 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.03.011
compatible with both the matrix and fibre, and thus could act as coupling agents [17,18]. For composites based on semicrystalline polymers, the morphological features like degree of crystallinity, spherulite size, lamella thickness and crystallite orientation play an important role in influencing the ultimate properties of the polymer matrix, and thus the composite. It was found that both morphology and crystallinity of thermoplastic polymer matrices deeply affected by fibre reinforcement [19–21]. In the case of thermoplastic composites, the growth of highly oriented transcrystalline layers at fibre–matrix interphase is thought to have a critical influence on improvement of the mechanical properties, due to better fibre–matrix adhesion [22]. The occurrence of transcrystallization depends on the type of fibre used and the crystallization temperature [23,24]. A transcrystalline layer can be formed at the fibre–matrix interface regions if the fibres with high nucleating ability were employed. However, there were not many studies devoted to investigate the crystallization behavior and microstructure of the natural fibre reinforced thermoplastic composites [23–28]. The mechanical properties of fibre-reinforced composites strongly depend on several factors such as fibre size, fibre loading, fibre dispersion, fibre orientation and fibre–matrix interfacial bond strength [29,30]. In many occasions, the effects of fibre reinforcement on the mechanical properties have been explained by comparing the experimental results with existing micromechanical models such as Parallel and Series, Hiesch, Cox, Halpin–Tasi, Modified Halpin–Tasi, Bowyer–Bader models, etc. [31–35]. It was
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
493
Table 1 Physical, chemical and mechanical properties of sisal fibres
Table 2 Formulation of the different HDPE/sisal composites
Fibre
Composites
Average diameter (m) Average length (mm) Density (g/cm3 ) Cellulose (%) Hemi-cellulose (%) Lignin (%) Tensile strength (MPa) Tensile modulus (MPa) Elongation (%)
Sisal 100 4.0 1.3–1.5 56.52 16.49 10.62 580 773 4.3
(95:5) HDPE/fibre (90:10) HDPE/fibre (85:15) HDPE/fibre (80:20) HDPE/fibre
Weight percentage (wt%) HDPE matrix
Ionomer-treated sisal fibre
Volume fraction of fibre, Vf (%)
95 90 85 80
5 10 15 20
3.22 6.44 9.66 12.88
2.4. Composites fabrication
concluded that the mechanical properties of fibre-reinforced composites strongly depend upon the fibre size distribution and fibre orientation distribution. However, a limited number of research works was dedicated to assess the viability of experimental data by using theoretical models for randomly oriented short fibre reinforced composites. The aim of this work is to investigate the effect of sisal fibre reinforcement on the crystallization kinetics and morphology of high-density polyethylene (HDPE). Surlyn ionomer was used as a coupling agent to improve the composites performance. The variations of tensile, flexural and impact properties as a function of fibre contents were studied. Moreover, a comparison between experimental and estimated Young’s modulus using various micromechanical models was presented for the HDPE/sisal composite.
2. Experimental 2.1. Materials The pure high-density polyethylene (M6805U: melt flow index = 0.5 g/10 min at 190 ◦ C and density = 0.968 g/cm3 at 23 ◦ C) was obtained from Haldia Petrochemical Ltd. India. The coupling agent used in this investigation was Surlyn® 1650 (density = 0.94 g/cm3 ; melt flow index = 1.5 g/10 min at 190 ◦ C; melting temperature = 97 ◦ C with a heating rate of 10 ◦ C/min; methacrylic acid co-monomer content = 3.7–4.0 mol% and 57 mol% acid was neutralized by sodium salt), procured from Du Pont USA. Sisal fibres (A. sisalana) were collected form local sources. Table 1 presents the physical, chemical and mechanical properties of sisal fibres. Sodium hydroxide (NaOH) and acetic acid used for chemical treatment of the fibres were obtained from E. Merck, India.
2.2. Washing and alkali treatment
Composites were prepared by the melt mixing of the ionomertreated short sisal fibres with HDPE in a HAAKE rotor mixture (Model Rheomix 600, Dreieich, Germany) at 120 ◦ C with roller blades and a mixing chamber with a 60 cm3 volumetric capacity. The process was carried out for 10 min at an optimum speed of 50–60 rpm, which depended on the quantity of the fibres by weight. Each batch contained various plastic to fibre weight ratio (95:5, 90:10, 85:15 and 80:20). The composite formulations are shown in Table 2. Each composite mixture was then homogenized in a tworoll mill (150E-400 Collins, Germany) at 130 ◦ C and compression molded with a Delta Malikson 100TY pressman (Mumbai, India) to produce composite sheets (3 ± 0.2-mm thick). Test specimens were prepared from these sheets according to ASTM (American Society for Testing and Materials) standard D with a counter-cut copy-milling machine (6490, Ceast, Italy) with calibrated templates. 2.5. Thermal analysis The thermal behavior of the neat HDPE and formulated HDPE/sisal composites was measured with a differential scanning calorimetric (DSC) analyzer (DSC-10Q, TA Instruments, New Castle, DE, USA) under nitrogen atmosphere. The samples were heated to 170 ◦ C with a heating rate of 10 ◦ C/min and kept at this temperature for 10 min to eliminate previous thermal history and then rapidly cooled (at 100 ◦ C/min) to crystallization temperature (Tc ), and maintained at Tc till the time necessary for complete crystallization of the polymer matrix. The heat evolved during the isothermal crystallization (Hc ) was recorded as a function of time (t), at different crystallization temperatures. After crystallization, the samples were heated to melting point at a rate of 10 ◦ C/min. The crystalline melting temperatures (Tm ) and degree of crystallinity (Xcr ) of the samples were determined from the endotherm. The enthalpy of crystallization of 100% crystalline polyethylene was taken as 290 J/g [36].
The short sisal fibres (length ∼4.0 mm) were thoroughly washed by aqueous detergent solution to remove dirt followed by cleaning with distilled water and dried in vacuum oven at 70 ◦ C for 1 day. The washed fibres were immersed in 5% aqueous NaOH solution for a period of 1 h at 40 ◦ C, and this was followed by washing with 0.1N acetic acid and distilled water. The alkali-treated fibres were than dried in vacuum oven at 70 ◦ C for 2 days to obtain mercerized fibres.
2.6. Mechanical properties
2.3. Coupling agent treatment
2.7. Morphological analysis
The mercerized sisal fibres were immersed in hot 5% Surlyn ionomer solution (in toluene) at 70 ◦ C for 30 min to obtain ionomertreated fibres.
The surface morphology of the composites was analyzed using a polarized optical microscopy (Leika Metalographic Aristomet model), which attached with automatic hot-stage thermal
Rectangular bar specimens for tensile and flexural tests were prepared according to ASTM D638 and ASTM D790, respectively. The measurements of the samples were conducted with an Instron 3366 tensile test machine (Grove City, PA, USA) at a crosshead speed of 2 mm/min. The Izod impact strength of notched samples (2 mm depth) was measured with a Davenport Izod impact tester as per ASTM D256 method.
494
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
Table 3 Melting and crystallization kinetic parameters of neat HDPE and HDPE/sisal composites obtained from DSC measurements Material HDPE/fibre
Thermal properties Tc a (◦ C)
nb
Kn c
t0.5 d (s)
Tm e (◦ C)
Hm f (J/g)
Eg (kJ/mole)
% Xcr h
100/0 100/0 100/0 100/0
116 119 121 124
2.06 2.19 2.43 2.62
0.05643 0.00255 0.00021 0.00003
374 868 1288 1405
126.6 127.1 129.8 131.4
82.94 86.42 89.03 90.48
85.8
28.6 29.8 30.7 31.2
95/5 95/5 95/5 95/5
116 119 121 124
2.32 2.51 2.74 2.93
0.14148 0.01014 0.00091 0.00008
254 628 726 945
124.7 125.5 127.6 130.4
86.42 90.48 91.35 96.86
108.3
29.8 31.2 30.5 33.4
90/10 90/10 90/10 90/10
116 119 121 124
2.28 2.79 2.83 2.88
0.46613 0.01588 0.00179 0.00017
232 511 610 855
123.7 124.4 126.4 128.7
92.22 95.41 97.15 100.63
127.6
31.8 33.9 34.5 34.7
85/15 85/15 85/15 85/15
116 119 121 124
2.68 2.73 2.96 3.01
0.52967 0.02548 0.00231 0.00016
208 458 485 642
123.3 123.8 125.5 127.9
98.6 102.66 107.01 109.04
134.8
33.0 33.4 36.9 37.6
80/20 80/20 80/20 80/20
116 119 121 124
2.85 2.93 3.25 3.51
0.72664 0.02734 0.00355 0.00033
153 376 422 541
122.8 124.5 125.3 127.6
108.87 110.35 117.79 116.49
151.5
36.5 37.0 39.5 39.0
a b c d e f g h
Crystallization temperature. Avrami exponent. Crystallization kinetic constant. Crystallization half time. Crystalline melting temperature. Heat of fusion. Crystallization activation energy. Crystallinity.
unit (Mettler FP-90). The composite samples were sandwiched between microscope glass slides and inserted into the hot stage unit. The samples were heated to 170 ◦ C at 10 ◦ C/min, held at this temperature for 10 min, then rapidly cooled to 116 ◦ C at 100 ◦ C/min and allowed to crystallize isothermally. Images were obtained using a Nikon digital camera connected to the microscopy.
2.8. Fourier transform infrared (FT-IR) spectroscopy Infrared spectroscopic analysis of the ionomer-treated and untreated sisal fibres was carried out using an FT-IR Thermo Nicolet (NEXUS 870) spectrophotometer at 25 ◦ C.
3. Results and discussion 3.1. Melting and crystallization behavior of HDPE/sisal composites The quantitative results of melting and crystallization behavior as determined from DSC measurements are summarized in Table 3. The crystalline melting temperature of HDPE matrix in the composites was recorded lower compared to neat HDPE. The Tm values in the composites decreased with increasing fibre content (Table 3). The occurrence of lower melting peak in the composites could be attributed to the fact that strong nucleation on fibre surfaces shortened the time required for HDPE crystallization, therefore, limiting the isothermal thickening of HDPE crystals and decreasing their melting temperature. For all examined samples, an increase
Fig. 1. Heat flow curves of (a) neat HDPE and (b) 80:20 HDPE/sisal composite during isothermal crystallization.
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
495
Fig. 2. Plots of relative crystallinity (Xt ) versus time (t) at various Tc obtained for (a) neat HDPE, (b) 90:10 HDPE/sisal composites and (c) 80:20 HDPE/sisal composites.
of the melting temperature (Tm ) with increasing Tc was observed. The overall crystallization rate (1/t0.5 ) at higher Tc was found to slow (Table 3), which might be caused the formation of thick HDPE crystals having higher melting temperature. The percent crystallinity of HDPE (%Xcr ) was significantly increased by fibre reinforcement, indicating strong heterogeneous nucleating ability of the sisal fibres. In addition, the nucleating effects of the fibre surfaces promoted the growth and formation of interfacial transcrystalline layer when the composite sample was cooled from the melt, as observed in the optical microscopy images [37,38]. The effect of sisal fibre reinforcement on the crystallization behavior and thermodynamics of HDPE matrix was evaluated by analyzing the isothermal crystallization kinetics of neat HDPE and its composites in the temperature range of 116–124 ◦ C. Fig. 1 shows the heat flow curves of 100% HDPE and 20% fibre filled HDPE/sisal composite during isothermal crystallization at various Tc . It was seen from Fig. 1 that at any crystallization temperature, the primary crystallization of HDPE matrix in the composites was almost completed before the crystallization of neat HDPE had reached to 50–70% conversion. The fast crystallization conversion in the composites could be attributed to the transcrystallization as a result of nucleating effects of fibre surfaces. The short induction time and high crystallization rate are supporting this hypothesis. The plots of relative crystallinity (Xt ) as function of time t at various crystallization temperatures (Tc = 116, 119, 121 and 124 ◦ C) are shown in Fig. 2. From the comparison among the crystallization isotherms at the same crystallization temperature it can be clearly pointed out that the crystallization rate of HDPE matrix in the composites was faster than that of neat HDPE, indicating the occurrence of
nucleating effects of sisal fibres in the composites. Fig. 3 illustrates the dependence of overall crystallization rate (Rc ), expressed as the reciprocal of crystallization half time (1/t0.5 ), on the crystallization temperature for all examined samples. The crystallization rate (Rc ) of composites was recorded higher compared to neat polymer at the same Tc , and it can be observed to increase with increasing fibre content (Fig. 3). One possible reason was that strong nucleation on
Fig. 3. Reciprocal of crystallization half time (t0.5 ) versus Tc for neat HDPE and different HDPE/sisal composites.
496
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
Fig. 4. Avrami plots of the (a) neat HDPE, (b) 90:10 HDPE/sisal composites and (c) 80:20 HDPE/sisal composites at various Tc .
fibre surfaces shortened the time required for HDPE crystallization. For all examined samples, the crystallization process was hindered with increasing Tc , as anticipated on the basis of the nucleation theory [39]. 3.2. Isothermal crystallization kinetics The crystallization kinetics of the neat polymer and its sisal fibre composites are analyzed by using the Avrami model [40]: Xt = 1 − exp(−Kn t n )
(1)
The linearisation of Eq. (1) leads to Eq. (2): ln[− ln(1 − Xt )] = n ln t + ln Kn
Avrami exponent in the composites was varied around 3.0 implying a heterogeneous nucleation with three-dimensional growth of HDPE crystals. In contrast, the value of n was varied about 2.0 for the neat HDPE indicating two-dimensional diffusion controlled crystal growth. At any crystallization temperature, the crystallization rate constant of HDPE exhibited a significant increment as the fibre content increase (Table 3). The higher crystallization rate constant of the composites could be ascribed to the nucleating effect of the fibres. The crystallization rate constant Kn can be presented by Arrhenius equation: 1/n
Kn (2)
where Kn and n are the crystallization rate constant and Avrami exponent, respectively. The Avrami approach is the most applied tool to describe the phenomenology of the crystal growth. The Avrami exponent depends on growth geometry and nucleation type of the crystals. According to Eq. (2), the values of Kn and n were obtained from the intercept and slope of the linear plots of ln[−ln(1 − Xt )] versus ln(t). Fig. 4 demonstrates the Avrami plots of the neat HDPE, (90:10) HDPE/sisal and (80:20) HDPE/sisal composites at various crystallization temperatures. For all examined samples, the values of Avrami parameters (n and Kn ) in the temperature range 116–124 ◦ C are complied in Table 3. As observed in Table 3 that n increased with increasing Tc , and it recorded slightly higher for the HDPE/sisal composites compared to neat HDPE. The
= K0 exp
−E RTc
(3)
where K0 , R and E are the pre-exponential factor, universal gas constant and activation energy, respectively. The value of activation energy for the crystallization process was calculated from the slopes of the linear plots of (1/n) ln Kn versus 1/Tc . The values of E of neat HDPE and composites are summarized in Table 3. The activation energy of HDPE crystallization was found to be significantly enhanced by the addition of short sisal fibre to HDPE matrix. The crystallization activation energy is the sum of the transport activation energy and the nucleation activation energy. As mentioned above, the nucleation effect of sisal fibres leads to enhance the crystallization of HDPE in the composites. Hence, the higher crystallization activation energy of the composite might be originated from their greater transport activation energy because the mobility of HDPE chain segments was strongly restricted
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
497
Fig. 5. Polarized optical micrographs of the isothermally crystallized 80:20 HDPE/sisal composites (Tc = 116 ◦ C).
Table 4 Mechanical properties of HDPE/sisal composites at different fibre content Sample
Mechanical properties Tensile strength (MPa)
Pure HDPE (95:5) HDPE/fibre (90:10) HDPE/fibre (85:15) HDPE/fibre (80:20) HDPE/fibre
26.0 28.5 30.7 32.0 32.8
± ± ± ± ±
0.5 1.9 1.8 2.3 2.8
Tensile modulus (MPa) 860 952 1086 1107 1219
± ± ± ± ±
0.3 2.3 3.7 3.3 3.4
by the fibres and consequently hindered the crystallization process. 3.3. Optical microscopy study As observed in the polarized optical microscopic analysis, the growth and morphology of the HDPE crystals were significantly influenced by the addition of sisal fibres. Fig. 5 shows the optical photomicrographs of the isothermally crystallized (80:20) HDPE/sisal composites (Tc = 116 ◦ C). As shown in the photomicrographs, the nucleation of HDPE spherulites as well as regular transcrystalline growth occurred along the fibre–matrix interfaces. However, the formation and growth of transcrystallinity in a composite depend upon the performances of fibre–matrix interfaces [41]. As seen in Fig. 5, the crystalline morphology at the fibre–matrix interface does not look like the morphology of the crystals in the bulk of the matrix. In this situation, the transcrystalline layer was grown isothermally at the fibre–matrix interface while the bulk of the matrix still remained in the molten form, thus the growth transcrystallinity was practically undisturbed. The thickness of the transcrystalline layer was roughly constant of about 30 m. The fibre–matrix interface region was made of needle-like compact crystals propagated vertically on the surface of fibre. In earlier investigations, the needle-like crystals have also been identified on polyethylene with a certain molecular weight range under isothermal conditions [42,43]. 3.4. Mechanical properties The properties of matrix and fibre are very important aspects in achieving good mechanical properties of the composites. The melt strength and processability are the most sensitive parameters to matrix polymer, whereas the modulus is dependent on fibre properties [44]. The mechanical properties (tensile strength, tensile modulus, flexural strength, flexural modulus, and impact strength) of neat HDPE and its formulated sisal fibre composites are summarized in Table 4. The composites have shown scattered response to mechanical properties, as presented in Table 4, because of diffi-
Flexural strength (MPa) 37.0 40.7 43.2 45.0 46.7
± ± ± ± ±
0.8 1.5 1.2 1.9 1.9
Flexural modulus (MPa) 928 1131 1194 1354 1468
± ± ± ± ±
0.3 2.7 3.7 2.9 3.5
Impact strength (J/m) 37.0 20.3 17.7 15.2 14.3
± ± ± ± ±
0.16 0.19 0.24 0.22 0.26
culty in proper dispersion of short sisal fibres and random fibre alignment in the matrix. However, all the composites exhibited improved mechanical properties compared to neat HDPE. When the fibre content raised up to 20%, the tensile strength, tensile modulus, flexural strength and flexural modulus of the HDPE/sisal composite increased about 27, 41, 26, and 58%, respectively. The strong interfacial adhesion between HDPE and ionomer-treated fibres lead to improve the stress transfer from the matrix to the fibres, which caused better mechanical performance. At lower fibre content (<10%) the dispersion of fibres was very poor that resulted poor stress transfer, whereas a strong tendency of the fibre–fibre interaction at higher fibre content (>15%) significantly enhanced the stress transfer. A hypothetical model of the interface between ionomer-treated sisal fibre and HDPE is shown in Scheme 1. The PE chain of ionomer diffuses into the HDPE matrix through interchain entanglements. On the other hand, the acrylic groups of ionomer molecules form ester bonds with the hydroxyl groups of the fibres. FT-IR spectra of ionomer-treated and untreated sisal fibres are reported in Fig. 6. The spectra of ionomer-treated sisal fibres show a significant absorption peak in the carbonyl region (1715 cm−1 ), which is associated with the stretching vibration of methacrylic ester bond.
3.5. Modelling of mechanical property The elastic properties of the HDPE/sisal composites can be theoretically determined or derived by using a variety of micromechanical composite models. These models are derived based on the properties of the individual components of the composite and their arrangement [45]. Thus, a micromechanical model might be used to find the best combination of constituent materials to satisfy material design considerations. In this study, four different micromechanical composite models have been compared with experimental results in order to assess the influences of fibre dispersion on the Young’s modulus of HDPE/sisal composites. The micromechanical models of the Young’s modulus for randomly distributed short fibre-based polymer composites are as follows.
498
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
Scheme 1. Hypothetical model of fibre modification with Surlyn ionomer and reaction with HDPE matrix.
3.5.1. Voigt–Reuss model The Voigt–Reuss model [46,47] has been used to estimate the Young’s modulus for a lamina of short random fibre composites. This model is independent on the geometry of reinforcement. It can be expressed as: εVR =
3 5 [V ε + (1 − Vf )εm ] + 8 f f 8
εf εm εf (1 − Vf ) + εm Vf
(4)
where εf and εm are the Young’s moduli of fibre and matrix, respectively. Vf is the volume fraction of fibre in the composites, calculated using the following equation: Vf =
Wf m Wf f + (1 − Wf )f
3.5.2. Halpin–Tsai model The semi-empirical equation developed by Halpin and Tsai [48] has been widely used for predicting the elastic properties of a lamina of short random fibre composites. The following form of the Halpin and Tsai equation was used to determine the Young’s modulus of the composites: εHT =
8
1 − L Vf
+
5 1 + 2T Vf εm 8 1 − T Vf
(6)
where lf and df are the average length and diameter of the fibres. In Eq. (6) the parameters L and T are given as:
(5)
where Wf , f and m denote the weight fraction of fibres in the composite, the density of fibre and the density of matrix, respectively.
3 1 + 2(l /d ) (1 − V ) f f L f
L =
(εf /εm ) − 1 (εf /εm ) + 2(lf /df )
(7)
T =
(εf /εm ) − 1 (εf /εm ) + 2
(8)
3.5.3. Modified Rule of Mixtures The Rule of Mixtures can be used to determine the Young’s modulus for three-dimensional random short fibre composites [49,50]. This model works extremely well for aligned continuous fibre composites. The Young’s modulus was determined by using this equation: εMRM =
1 V ε + (1 − Vf )εm 5 f f
(9)
3.5.4. Cox model This model introduces the influence of the reinforcement aspect ratio into the expression of Rule of Mixtures [51]. The Young’s modulus of three-dimensional random short fibre composites was evaluated using this equation:
Fig. 6. FT-IR spectra of ionomer-treated and untreated sisal fibres.
εCox =
1 L εf Vf + (1 − Vf )εm 5
(10)
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
499
Table 5 Experimental and predicted Young’s modulus for neat HDPE, sisal fibre and HDPE/sisal composites Sample
Young’s modulus, ε (GPa) Experimental
Pure HDPE Sisal fibre (95:5) HDPE/fibre (90:10) HDPE/fibre (85:15) HDPE/fibre (80:20) HDPE/fibre
1.195 12.40 2.002 2.425 2.975 3.124
± ± ± ± ± ±
0.1 2.7 0.2 0.3 0.2 0.3
Voigt–Reuss model
Halpin–Tsai model
Modified mixture law
Cox model
– – 3.333 4.688 5.439 5.947
– – 3.003 4.115 4.652 5.277
– – 2.313 3.033 3.553 3.771
– – 2.207 2.544 3.037 3.281
indicating that the assumptions made to derive these theoretical equations were closely agreeing with experimental conditions. The relative deviation between experimental and theoretical results increased with increasing fibre content, but it was well within agreeable limits for the Modified Rule of Mixtures and the Cox model. This suggested that the Modified Rule of Mixtures and Cox model can be applied to predict Young’s modulus of HDPE/sisal composites. However, the best theoretical–experimental correlation was obtained with the Cox relationship because it takes into account the fibre aspect ratio. Apart from that, the modulus of composite depends on the homogeneity of fibre dispersion at a given volume fraction of sisal fibre. Obviously, the higher modulus was obtained for better fibre distribution. The composite Young’s modulus was nearly the matrix modulus for poor fibre dispersion. 4. Conclusions
Fig. 7. Comparison of experimental Young’s modulus values of HDPE/sisal composites with different micromechanical models.
The independent parameters of Eq. (10) are calculated using the following equations:
L = 1 − s=
tan(ˇs) ˇs
2lf 0.5df
ˇ2 =
2εm εf (1 + m ) ln(1/Vf )
(11)
(12) (13)
where m is the Poisson’s ratio of HDPE matrix. Generally, Young’s modulus of the composite material reflects the capability of both fibre and matrix to transfer the elastic deformation under small strain without interface fracture. Table 5 presents the predicted Young’s modulus values obtained from micromechanical models and the experimental data for all formulated samples. It can be seen that the Young’s modulus of the composites increased as the fibre volume fraction increase. A comparison between the experimental data and predicted Young’s modulus values is demonstrated in Fig. 7. The predicted Young’s modulus values were recorded higher than experimental values (Fig. 7). Fig. 7 shows a greater degree of divergence between the experimental Young’s modulus values and the predicted values using Voigt–Reuss and Halpin–Tsai models, because these models were developed for lamina composite (two-dimensional structure), not for three-dimensional composite. The differences of these two micromechanical models with the experimental Young’s modulus were recorded ∼65% for the Voigt–Reuss model and ∼50% for the Halpin–Tsai model. Although, the comparison between experimental Young’s modulus and those obtained from Modified Rule of Mixtures or Cox model demonstrated a close agreement within 10%,
The present investigation showed the usefulness of sisal fibre as a good reinforcing agent and Surlyn ionomer as an effective coupling agent for composite fabrication with polyethylene matrix. DSC analysis demonstrated that the presence and the concentration of short sisal fibres affect markedly the crystallization behavior of HDPE matrix. A noticeable decrease in the half time of HDPE crystallization was observed by the incorporation of short sisal fibre to HDPE. This could be attributed to the fact that the surfaces of sisal fibres act as nucleating sites for the crystallization of the matrix polymer, promoting the growth and formation of transcrystalline layer in the composites. The concentration of nucleating sites in the HDPE/sisal composites increased as the fibre content increase, which was reflected in the increase of the crystallinity of the composites. The Avrami exponent in the composite samples was varied around 3.0 implying three-dimensional growth of HDPE crystals, whereas the n-values of about 2.0 for the neat HDPE suggest two-dimensional diffusion controlled crystal growth. A hot-stage polarized optical microscopy was used to follow the development of transcrystalline layer along the fibre–matrix interface. The mechanical properties of HDPE/sisal composites were found to increase with increasing fibre content. This was ascribed to the reinforcing effect imparted by the ionomer-treated fibres, which allowed a uniform stress distribution from the polymer matrix to the dispersed fibre phase. A comparison between the experimental results and the theoretically predicted Young’s modulus has been presented. Modified Rule of Mixtures and Cox model give more close agreement to experimental data for HDPE/sisal composite. References [1] K.G. Satyanarayana, B.C. Pai, K. Sukumaran, S.G.K. Pillai, Handbook of Ceramics and Composites, Marcel Decker, New York, 1990, pp. 339–360. [2] K.G. Satyanarayana, K. Sukumaran, P.S. Mukherjee, C. Pavithran, S.G.K. Pillai, Natural Fibre-Polymer Composites, Haworth Press, England, 1990, pp. 117–136. ´ adi, ´ ´ G. Nagy, J. Mocz ´ o, ´ B. Pukanszky, ´ [3] L. Dany K. Renner, Z. Szabo, Polym. Adv. Technol. 17 (2006) 967–974. [4] S. Mohanty, S.K. Verma, S.K. Nayak, Compos. Sci. Technol. 66 (2006) 538–547.
500
A. Choudhury / Materials Science and Engineering A 491 (2008) 492–500
[5] I. Baroulaki, J.A. Mergos, G. Pappa, P.A. Tarantili, D. Economides, K. Magoulas, C.T. Dervos, Polym. Adv. Technol. 17 (2006) 954–966. [6] T. Li, N. Yan, Compos. A 38 (2007) 1–12. [7] C. Pavithran, P.S. Mukharjee, M. Brahmakumar, A.D. Damodaran, J. Mater. Sci. Lett. 6 (1987) 882–884. [8] K. Jayaraman, Compos. Sci. Technol. 63 (2003) 367–374. [9] X. Colom, F. Carrasco, P. Pages, J. Canavate, Compos. Sci. Technol. 63 (2003) 161–169. [10] P.A. Sreekumar, K. Joseph, G. Unnikrishnan, S. Thomas, Compos. Sci. Technol. 67 (2007) 453–461. [11] S. Mohanty, S.K. Nayak, Mater. Sci. Engg. A443 (2007) 202–208. [12] N. Chand, D. Jain, Compos. A 36 (2005) 594–602. [13] M. Idicula, A. Boudenne, L. Umadevi, L. Ibos, Y. Candau, S. Thomas, Compos. Sci. Technol. 66 (2006) 2719–2725. [14] M. Brahmakumar, C. Pavithran, R.M. Pillai, Compos. Sci. Technol. 65 (2005) 563–569. [15] J.M. Park, S.T. Quang, B.S. Hwang, D.V.K. Lawrence, Compos. Sci. Technol. 66 (2006) 2686–2699. [16] L.M. Arzondo, C.J. Perez, J.M. Carella, Polym. Engg. Sci. 45 (2005) 613–621. [17] H. Akimoto, T. Kanazawa, M. Yamada, S. Matsuda, G.O. Shonaike, A. Murakami, J. Appl. Polym. Sci. 81 (2001) 1712–1720. [18] E.R. Sadiku, R.D. Sanderson, Macromol. Mater. Engg. 286 (2001) 472–479. [19] M. Arroyo, R. Zitzumbo, F. Avalos, Polymer 41 (2000) 6351–6359. [20] E. Piorkowska, Macromol. Symp. 169 (2001) 143–148. [21] H. Quan, Z.-M. Li, M.-B. Yang, R. Huang, Compos. Sci. Technol. 65 (2005) 999–1021. [22] H. Schonhorn, F.W. Ryan, J. Polym. Sci. 6 (1968) 231–240. [23] C. Wang, L.M. Hwang, J. Polym. Sci. B 34 (1996) 1435–1442. [24] C. Wang, C.R. Liu, Polymer 38 (1997) 4715–4718. [25] P.V. Joseph, K. Joseph, S. Thomas, C.K.S. Pillai, V.S. Prasad, G. Groeninckx, M. Sarkissova, Compos. A 34 (2003) 253–266. [26] P.V. Joseph, K. Joseph, S. Thomas, Compos. Sci. Technol. 59 (1999) 1625–1640. [27] J.L. Thomason, A.A. Van Rooyen, J. Mater. Sci. 27 (1992) 889–896.
[28] M. Pracella, D. Chionna, I. Anguillesi, Z. Kulinski, E. Piorkowska, Compos. Sci. Technol. 66 (2006) 2218–2230. [29] A.K. Rana, A. Mandal, S. Bandyopadhyay, Compos. Sci. Technol. 63 (2003) 801–806. [30] A.P. Kumar, R.P. Singh, B.D. Sarwade, Mater. Chem. Phys. 92 (2005) 458–469. [31] A.G. Facca, M.T. Kortschot, N. Yan, Compos. Sci. Technol. 67 (2007) 2454– 2466. [32] R.G. de Villoria, A. Miravete, Acta Mater. 55 (2007) 3025–3031. ¨ [33] T.T.L. Doan, S.L. Gao, E. Mader, Compos. Sci. Technol. 66 (2006) 952–963. [34] G. Kalaprasad, K. Joseph, S. Thomas, J. Mater. Sci. 32 (1997) 4261–4267. [35] G. Kalaprasad, K. Joseph, S. Thomas, J. Compos. Mater. 31 (1997) 509–527. [36] J. Brandup, E.H. Immergut, Polymer Handbook, Wiley, New York, 1988 (chapter 8). [37] Y. Huang, J. Petermann, Polym. Bull. 36 (1996) 517–524. [38] A. Amash, P. Zugenmaier, Polymer 41 (2000) 1589–1596. [39] J.D. Hoffman, G.T. Davis, J.I. Lauritzen, The Rate of Crystallization of Linear Polymers with Chain Folding, Crystalline and Noncrystalline Solids, Plenum Press, New York, 1976, p. 497. [40] M. Avrami, J. Chem. Phys. 7 (1939) 1103–1112. [41] S. Incardona, C. Migliaresi, D. Wagner, A. Gilbert, G. Marom, Compos. Sci. Technol. 47 (1993) 43–50. [42] L. Mandelkern, M. Glotin, R.A. Benson, Macromolecules 14 (1981) 22–34. [43] J. Maxfield, L. Mandelkern, Macromolecules 10 (1977) 1141–1153. [44] D. Hull, An Introduction to Composite Materials, 2nd ed., Cambridge University Press, Cambridge, 1996, pp. 326–345. [45] R.M. Jones, Mechanics of Composite Material, Hemisphere Publishing Corporation, New York, 1975, pp. 86–104. [46] W. Voigt, Ann. Phys. 38 (1889) 573–587. [47] A. Reuss, Z. Angrew. Math. Mech. 9 (1929) 49–58. [48] J.C. Halpin, J. Compos. Mater. 3 (1969) 732–734. [49] H. Fukuda, T.W. Chou, J. Mater. Sci. 16 (1981) 1088–1096. [50] H. Fukuda, T.W. Chou, J. Mater. Sci. 17 (1982) 1003–1011. [51] H.L. Cox, Brit. J. Appl. Phys. 3 (1952) 72–79.