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On the post-mortem fracture surface morphology of short fiber reinforced thermoplastics
Composites: Part A 36 (2005) 987–994 www.elsevier.com/locate/compositesa
On the post-mortem fracture surface morphology of short fiber reinforced thermoplastics S.Y. Fua,b,c,*, B. Laukeb, Y.H. Zhanga, Y.-W. Maic a
Technical Institute of Physics and Chemistry, Chinese Academy of Sciences, P.O. Box 2711, Beijing 100080, China b Leibniz Institute of Polymer Research, Hohe Street 6, D-01069 Dresden, Germany c Centre for Advanced Materials and Technology, School of Aerospace, Mechanical and Mechatronic Engineering (J07), The University of Sydney, Sydney NSW 2006, Australia Received 5 March 2004; revised 18 November 2004; accepted 18 November 2004
Abstract The fracture surface morphology of short fiber reinforced thermoplastics (SFRTs) has often been used to assess qualitatively the degree of fiber–matrix interfacial adhesion. Mechanical properties such as tensile strength, fracture toughness and failure strain, etc. are then correlated with the morphology. Fracture surfaces showing fibers surrounded by a large amount of matrix material is commonly regarded as indication of strong fiber–matrix interfacial adhesion while smooth fibers are characteristic of weak interfacial adhesion. Many experimental results of SFRTs have been so interpreted. However, it is shown in this paper that strictly speaking, such interpretations are generally incorrect. Moreover, the amount of matrix material does not provide a quantitative measure of the adhesion. Correct implication of the morphology of fracture surfaces is clarified. Short glass fiber reinforced polyamide 6,6/polypropylene (PA 6,6/PP) blends toughened by rubber are employed as examples for SFRTs since the PA 6,6/PP blend system by changing PA 6,6 concentration in the matrix blend represents a wide range of matrix materials. It is demonstrated that the fracture surface morphology of such composites is dependent on both fiber–matrix interfacial adhesion strength and matrix shear yield strength. Consequently, tensile failure strain is well correlated with the post-mortem fracture surface morphology of these SFRTs. q 2004 Elsevier Ltd. All rights reserved. Keywords: A. Glass fibres; B. Interface/interphase; Thermoplastics
1. Introduction Short fiber reinforced thermoplastics (SFRTs) are very attractive because of their ease of fabrication, economy and superior mechanical properties [1–6]. Characterization of the morphology of the post-mortem fracture surfaces of SFRTs is often conducted with scanning electron microscopy (SEM). These morphologies provide direct evidence to study the strengthening and toughening mechanisms of these composites. It has been reported that fracture surfaces of SFRTs containing fibers surrounded by a substantial amount of matrix material imply strong * Corresponding author. Address: Technical Institute of Physics and Chemistry, Chinese Academy of Sciences, P.O. Box 2711, Beijing 100080, China. Tel.: C86 10 62659040; fax: C86 10 62564049. E-mail address: [email protected] (S.Y. Fu). 1359-835X/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.compositesa.2004.11.005
interfacial adhesion between fibers and matrix [7–9]. Conversely, smooth fiber surfaces signify weak interfacial adhesion between fibers and matrix [7–10]. The morphology of the fracture surfaces has often been used to correlate with the mechanical properties. Alsewailem concluded in his PhD project [7], from observation of the fracture surfaces with fibers surrounded by a large amount of matrix material, that addition of rubber to nylon 66 led to strong glass fiber–matrix interfacial adhesion, which in turn increased the impact strength. Nair et al. [8] attempted to link the fracture toughness of glass fiber reinforced rubbertoughened nylon 6,6 composites to their post-mortem fracture morphologies, such that clean pullout fibers on the fracture surfaces implied weak fiber–matrix interfaces while profound matrix-adhered fibers denoted strong fiber– matrix interfaces. It was then claimed that when the fiber– matrix interface was strong, the composites were tougher
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than the unreinforced polymer matrices; when the fiber– matrix interface was weak, the interface had little effect on the toughness. On the other hand, Stamhuis [10] compared the fracture surface morphologies of PP/GF/EPDM and PP/ GF/SBS composites and found no differences due to the two ealstomers (EPDM and SBS). However, better impact behavior was found for the PP/GF/EPDM composite. Further evidence is provided by Cho and Paul [11] who studied the mechanical properties and morphology of glass fiber reinforced ABS toughened nylon 6 composites. The ABS material used was a high rubber, emulsion-made product. Smooth fiber surfaces were observed for the composite without ABS. When the matrix was blended with ABS, there appeared to be some matrix material adhered to the fiber surfaces and the amount seemed to vary directly with rubber content. If the fracture surface morphology—interfacial adhesion—composite mechanical properties correlation is correct, the composite strength would increase with rubber content due to the enhanced interfacial adhesion. However, it was reported [11] that the composite strength actually decreased with increasing ABS content. Therefore, there is concrete evidence from previous studies to show the limitation of linking mechanical properties to post-mortem fracture surface morphology. The purpose of the present work is to clarify the implication of the morphology of the fracture surfaces of SFRTs. Tensile failure strain is then correlated with the post-mortem fracture surface morphology. Polyamide 6,6 (PA 6,6)/polypropylene (PP) blends were used as matrices since the PA 6,6/PP blend system by changing the PA 6,6 concentration in the matrix represented a wide range of matrix materials. A mixture of 10 wt% styrene–ethylene– butylene–styrene (SEBS) and 10 wt% maleic anhydride (MA) grafted SEBS (SEBS-g-MA), which was proven in earlier studies to be effective both as a compatabilizer and a toughener for PA 6,6/PP blends [12,13], was used to compatabilise the matrices. Short glass fibers (SGF) purchased from Owens Corning (123D) were used to reinforce the PA 6,6/PP blends.
2. Experimental details 2.1. Materials and processing Polyamide 6,6 (PA 6,6) resin (Vydyne 21) was obtained from Monsanto Chemical Company (USA) and polypropylene (PP) (Propathene GSE52) was supplied by ICI, Australia. The triblock copolymer SEBS (Kraton G1652) and its maleated version (Kraton F1901X) were bought from Shell Chemical. 123D short glass fibers of a length of 4 mm and an average diameter of 10 mm were purchased from Owens Corning. PA compatible sizing has been made by the producer on the surfaces of 123D glass fibers for reinforcing PA and their blends where hydrolysis resistance is not an end use requirement.
Polymer materials were first dry blended simultaneously at various weight ratios (100/0, 75/25, 50/50, 25/75 and 0/100) of PA 6,6/PP, 80 wt% PA 6,6/PP blend plus a mixture of 10 wt% non-maleated-SEBS and 10 wt% maleated SEBS copolymer. Then, short glass fibers were pre-compounded together with the dry blended polymer matrices at a weight ratio of 40/60 glass/polymer. Compounding of the materials was carried out using a high shear-rate twin-screw extruder (Werner and Pfleiderer ZSK 30) at a temperature range of 260–280 8C. Materials were injection molded into 3 mm thick tensile bars (ASTM D638 type I) by an injection molding machine (BOY 22S Dipronic). All materials were dried at 80 8C in an oven over night prior to compounding and injection molding. Samples from polymer materials were also prepared under the similar processing conditions so that the tensile data from these samples can be used as that of the matrix blends in the composite samples. 2.2. Tensile testing and post-mortem fracture morphology Tensile tests were conducted at room temperature at a crosshead rate of 5 mm/min using an Instron Model 5567 computer-controlled testing machine. Six samples for each composition were tested immediately after preparation of the samples. The samples were not humidity conditioned before testing. Scanning electron microscopy (SEM) examination was conducted on the fracture surfaces of the samples. Before SEM examination, samples were gold sputtered by a JEOL JFC-1100E surface coating machine and viewed with a JEOL JSM-820 SEM. 2.3. Fiber dispersion, mean fiber length and critical fiber length Optical micrographs of the fiber dispersion across the sections of the samples perpendicular to the flow direction were taken after the cross-section were polished. Fiber lengths were determined by burning off the polymer matrix in a furnace at 550 8C for 1 h. The fibers were then dispersed with water on glass slides so that their lengths could be determined microscopically after the water evaporated. About 600–1000 fibers were measured for each composition to get the mean fiber lengths. Critical fiber lengths were approximately obtained by doubling the average of the maximum pull-out fiber lengths from the fracture surfaces of the samples for each composition since the maximum fiber pull-out length should be equivalent to half of the critical length in principle. Pull-out lengths of over 30 longest pulled-out fibers (shown in Fig. 1) were measured and the 10 maximum fiber pull-out lengths were averaged for each composition as half of the critical fiber length. By the way, another method can also be employed to estimate the critical fiber length. First, all the pullout lengths of the fibers, which vary from 0 to half critical length, are
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Fig. 1. SEM micrographs of fiber pullout at the fracture surfaces of 40 wt% SGF reinforced PA 6,6/PP composites. (A) 100% PA 6,6, (B) 75% PA 6,6, (C) 50% PA 6,6, (D) 25% PA 6,6 and (E) 0% PA 6,6 in the matrix.
estimated, then the average of all the measurements is multiplied by 4 to get the critical fiber length. Similar results can be obtained.
3. Results and discussion 3.1. Effect of PA 6,6 concentration in the matrix blend on interface adhesion strength Fig. 2 shows the tensile strength of SGF reinforced PA 6,6/PP blends. It can be seen that the composite tensile strength increases with increasing the PA 6,6 concentration in the matrix system. When the PA 6,6 concentration is equal to or higher than 75 wt% in the matrix blend, the composite strength is much higher than for other PA 6,6 concentrations. Table 1 indicates that the mean fiber length decreases with increasing the PA 6,6 concentration in the matrix system. The increase in PA 6,6 concentration in
Fig. 2. Effect of the PA 6,6 concentration in the PA 6,6/PP matrix system on tensile strength of 40 wt% SGF reinforced PA 6,6/PP composites.
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Table 1 Various quantities versus PA 6,6 concentration in the matrix blend PA 6,6/matrix (PA 6,6CPP) (%)
Fiber volume fraction (%)a
Mean fiber length (mm)
Critical fiber length (mm)
Interfacial adhesion strength (MPa)
Matrix tensile yield strength (MPa)
Matrix shear yield strength (MPa)
100 75 50 25 0
0.220 0.213 0.205 0.198 0.190
182.3 209.0 268.6 324.4 332.2
260.2 275.8 430.1 650.0 950.6
38.43 36.26 23.25 15.38 10.52
40.05 32.10 20.93 14.98 19.69
20.03 16.05 10.47 7.49 9.85
a
The densities of PA 6,6, PP, SEBS and glass fibers used for the computation of fiber volume fraction are, respectively, 1.13, 0.91, 0.88 and 2.55 g/cm3.
the matrix blend corresponds to an increase in fiber volume fraction as shown in Table 1. It is well known that the viscosity of polymers can be significantly increased by the addition of fibers. Therefore, the increase of fiber volume fraction would lead to lower average fiber length due to higher viscous forces exerting on fibers. On the other hand, as fiber volume fraction increases, fiber–fiber interaction would also cause more damage to fiber length [14]. It was
shown previously [15,16] that the tensile strength of short fiber reinforced thermoplastic composites increased with both fiber–matrix interfacial strength and mean fiber length since fiber orientation and dispersion are approximately the same for the investigated cases (please see Fig. 3 in which fibers are uniformly dispersed and preferentially aligned along the flow direction, namely sample axis direction in all the investigated samples). It has been shown in Table 1 that
Fig. 3. Optical micrographs of the cross-sections of the composites. (A) 100% PA 6,6, (B) 75% PA 6,6, (C) 50% PA 6,6, (D) 25% PA 6,6 and (E) 0% PA 6,6 in the PA 6,6/PP matrix system.
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mean fiber length decreases with increasing the PA 6,6 concentration in the matrix system, this would lead to reduction in composite tensile strength. Nonetheless, it can be observed that the composite strength increases with the increase of the PA 6,6 concentration. It can thus be inferred that the fiber–matrix interfacial adhesion strength must increase with the PA 6,6 concentration in the matrix system. This is further confirmed below. It is shown that the critical fiber length decreases as the fiber–matrix interfacial adhesion strength increases [15,16]. The relationship between critical fiber length (lf) and interfacial adhesion strength (ti) is given by the wellknown equation below [17]: lc Z rf sfu =ti
(1)
where rf is fiber radius and sfu fiber fracture strength. The fracture strength of fibers used here is 2 GPa [18]. Some fractured fibers can be observed at the fracture surfaces of the composites (Fig. 4), indicating that longest fibers at
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the fracture surfaces should have approximately critical fiber lengths because longer fibers would be fractured. Thus, the approximation of the critical fiber lengths can be obtained from Fig. 1 for various matrix materials and then, the interfacial adhesion strength can be estimated according to Eq. (1). The results are given in Table 1. Indeed, it is noted that the critical fiber length increases with the decrease of the PA 6,6 concentration in the matrix system and hence, the interfacial adhesion strength increases with increasing the PA 6,6 concentration. 3.2. The post-mortem morphology of the pullout fibers at the fracture surfaces of SFRTs It has been shown above that the interface adhesion strength increases with the increase of the PA 6,6 concentration in the matrix system. As a result, it can be conjectured that the amount of the matrix material surrounding the fibers standing proud at the fracture surfaces
Fig. 4. SEM micrographs of the fracture surfaces of the composites. (A) 100% PA 6,6, (B) 75% PA 6,6, (C) 50% PA 6,6, (D) 25% PA 6,6 and (E) 0% PA 6,6 in the PA 6,6/PP matrix system.
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Fig. 5. The morphologies of the fracture surfaces of 40 wt% SGF reinforced PA 6,6/PP composites. (A) 100% PA 6,6, (B) 75% PA 6,6, (C) 50% PA 6,6, (D) 25% PA 6,6 and (E) 0% PA 6,6 in the PA 6,6/PP matrix system.
should also increase with the increase of the PA 6,6 concentration in the matrix system according to the common belief on the morphology of the pullout fibers. However, Fig. 5 shows that the medium PA 6,6 concentrations (75, 50 and 25%) in the matrix system correspond to greater amounts of the matrix materials surrounding the fibers rather than for the 100% PA 6,6 case. The micrographs in Fig. 5 are arbitrarily selected for the observation of the morphology of the fiber surfaces. Micrographs containing a number of fibers with the similar morphology of fiber surfaces are available in Fig. 4 with a lower magnification. Thus, Fig. 5 can be regarded to represent the observation of the fiber surfaces at the whole fracture surfaces. Consequently, the amount of the matrix material does not provide a quantitative measure of the adhesion and the fiber–matrix interface adhesion cannot be simply linked to the post-mortem morphology of the pullout fibers on the fracture surfaces. The traditional interpretations presented previously for the morphology of the fracture surfaces of SFRTs are thus incorrect.
Table 1 shows the values for the matrix tensile yield strength obtained from tensile testing of matrix materials. Based on Mohr’s theory, the matrix shear yield strength must be about 50% the matrix tensile yield strength and the results for the matrix shear yield strength are then given in Table 1. It can be seen that the interfacial adhesion strength is obviously higher than the matrix shear yield strength except for the case of the 0% PA 6,6 in the matrix system where the interface adhesion strength is close to the matrix shear yield strength. It can then be reasonably inferred that the matrix shear yielding and then failure would occur before the interface failure except for the case of the 0% PA 6,6 in the matrix system. Thus fiber surfaces adhered with matrix materials have been observed except for the 0% PA 6,6 in the matrix system as shown in Fig. 5. For the SFRT containing no PA 6,6, fiber surfaces are quite smooth because interface debonding would take place before occurrence of matrix shear yielding due to the fact that the interface stress at fiber ends must be quite higher than
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Fig. 7. Stress–strain curves for 123D SGF reinforced PA 6,6/PP composites. Fig. 6. Effect of the PA 6,6 concentration in the matrix system on tensile failure strain of 40 wt% SGF reinforced PA 6,6/PP composites.
the average interface shear stress [19,20] while the interface adhesion strength is very close to the matrix shear yield strength. Consequently, the post-mortem morphology of the fracture surfaces is determined by both interface adhesion strength and matrix shear yield strength. 3.3. Correlation of the post-mortem morphology of the pullout fibers with composite failure strain Comparison of Figs. 5 and 6 shows that the changing trend of the morphology is the same as that for the failure strain with the PA 6,6 concentration in the matrix system. At the one extreme of 0% PA 6,6 (namely 100% PP) where interface adhesion is weakest and fiber surfaces are smooth, the tensile failure strain is the lowest. At another extreme of 100% PA 6,6 where interface adhesion is strongest but there is some amount of matrix adhering to fiber surfaces, the tensile failure strain is higher than for the 0/100 PA 6,6/PP SFRP but lower than for other composites. While between the two extremes, substantial amounts of matrix material surrounding fiber surfaces are observed, the failure strains are higher than for the two extremes. From the tensile curves of the composites shown in Fig. 7, it can be seen that the tensile stress–strain curves show linear deformation at lower stresses and non-linear deformation at higher stresses as having been shown previously [4]. The deviation of the curves from linear to non-linear deformation corresponds to either interfacial debonding or matrix shear yielding deformation. For the PA 6,6 based composite (100% PA 6,6 in the matrix system), the tensile curve deviates from linear to non-linear behavior at a relatively higher stress than for other composites but the deformation from occurrence of non-linear deformation to composite fracture is relatively small (namely a small matrix shear yielding deformation) when compared with other composites, bringing about a relatively small failure strain and thus a small amount of matrix material adhering to fiber surfaces because of the small matrix shear yielding
deformation. For the SFRT containing no PA 6,6 with a low interfacial adhesion strength, interfacial debonding would occur at a relatively low stress and the non-linear deformation does not last long to composite fracture due to the weak interface adhesion, leading to the lowest failure strain of the composite. This is because interface debonding takes place before matrix shear yielding, no or little matrix material is adhered to fiber surfaces and thus smooth fiber surfaces are observed. For the SFRT containing 75, 50 and 25% PA 6,6 in the matrix system, the non-linear deformation is relatively large, resulting in relatively large failure strains (corresponding to relatively large matrix deformation and thus relatively large amounts of matrix material adhering to fiber surfaces). Hence, in summary, the tensile failure strain can be well linked to the post-mortem morphology of the pullout fibers on the fracture surfaces. 4. Conclusions The implications of the post-mortem morphology of the pullout fibers on the fracture surfaces of SGF reinforced PA 6,6/PP blends as an example of SFRTs are discussed. The results have shown that the fiber–matrix interface adhesion cannot be simply linked to the post-mortem morphology of the pullout fibers on the fracture surfaces. The traditional interpretations presented previously for the morphology of the fracture surfaces of SFRTs are thus incorrect. The observation of the fibers surrounded by a great amount of matrix material at the fracture surfaces does not certainly denote a strong fiber–matrix interfacial adhesion. Moreover, it is displayed that the post-mortem morphology of the pullout fibers on the fracture surfaces is determined by both interface adhesion strength and matrix shear yield strength. Finally, it is observed that the tensile failure strain can be well linked to the post-mortem morphology. Acknowledgments This work was funded by the Talent’s Program, Chinese Academy of Sciences and the Australian Research Council
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on ‘Short Glass Fiber Reinforced Toughened Polymer Blends’. YWM was supported by an Australian Federation Fellowship. SYF is grateful for the support of the Alexander von Humboldt Foundation and the Australian Research Council.
References [1] Ramsteiner F, Theysohn R. Tensile and impact strengths of unidirectional, short fiber-reinforced thermoplastics. Composites 1979;10:111–9. [2] Fu SY, Lauke B. Fracture resistance of unfilled and calcite-filled ABS composites reinforced by short glass fibers (SGF) under impact load. Composites, Part A 1998;29:631–41. [3] Fu SY, Lauke B. Characterization of tensile behavior of hybrid short glass fibre/calcite particle/ABS composites. Composites, Part A 1998; 29:575–83. [4] Fu SY, Lauke B, Ma¨der E, Yue CY, Hu X. Tensile properties of shortglass-fiber- and short-carbon-fiber-reinforced polypropylene composites. Composites, Part A 2000;31:1117–25. [5] Friedrich K. Microstructural efficiency and fracture toughness of short fiber/thermoplastic matrix composites. Compos Sci Technol 1985;22: 43–74. [6] Jancar J. Yielding and impact behavior of pp/sgf/epr ternary composites with controlled morphology. J Mater Sci 1996;31:3983–7. [7] Alsewailem FD. Rubber toughening of glass-fiber-reinforced nylon 66. PhD thesis, Western Virginia University; 2002. p. 117–30. [8] Nair SV, Subramanian A, Goettler LA. Fracture resistance of polyblends and polyblend matrix composites. Part III. Role of rubber type and location in nylon 6,6/SAN composites. J Mater Sci 1998;33: 3455–64. [9] Hargarter N, Friedrich K, Cartsman P. Mechanical properties of glass fiber/talc/polybutylene–terephthalate composites as processed by the radlite technique. Compos Sci Technol 1993;46:229–44.
[10] Stamhuis JE. Mechanical properties and morphology of polypropylene composites. III. Short glass fiber reinforced elastomer modified polypropylene. Polym Compos 1998;4:280–4. [11] Cho JW, Paul DR. Glass fiber-reinforced polyamide composites toughened with ABS and EPR-g-MA. J Appl Polym Sci 2001;80: 484–97. [12] Wong SC, Mai Y-W. Effect of rubber functionality on microstructures and fracture toughness of impact-modified nylon 6,6/polypropylene blends. Part I. Structure–property relationships. Polymer 1999;40: 1553–66. [13] Wong SC, Mai Y-W. Effect of rubber functionality on microstructures and fracture toughness of impact-modified nylon 6,6/polypropylene blends. Part II—toughening mechanisms. Polymer 2000; 41:5471–83. [14] Fu SY, Hu X, Yue CY. Effects of fiber length and orientation distributions on the mechanical properties of short-fiber-reinforced polymers—a review. Mater Sci Res Int 1999;5:74–83. [15] Fu SY, Lauke B. Effects of fiber length and fiber orientation distributions on the tensile strength of the tensile strength of short-fiber-reinforced polymers. Compos Sci Technol 1996;56: 1179–90. [16] Templeton PA. Strength predictions of injection molding compounds. J Reinf Plast Compos 1990;9:210–25. [17] Fu SY, Lauke B. The fiber pull-out energy of misaligned short fiber composites. J Mater Sci 1997;32:1985–93. [18] Thomason JL. The influence of fibre length and concentration on the properties of glass fibre reinforced polypropylene: 5. Injection moulded long and short fibre PP. Composites, Part A 2002;33: 1641–52. [19] Fu SY, Zhou BL, Chen X, He GH, Lung CW. Some further considerations of the theory of fibre debonding and pull-out from an elastic matrix, I. Constant interfacial frictional shear stress. Composites 1993;24:5–11. [20] Fu SY, Yue CY, Hu X, Mai YW. Analyses of the micromechanics of the stress transfer in single- and multi-fiber pull-out tests. Compos Sci Technol 2000;60:569–79.
On the post-mortem fracture surface morphology of short fiber reinforced thermoplastics S.Y. Fua,b,c,*, B. Laukeb, Y.H. Zhanga, Y.-W. Maic a
Technical Institute of Physics and Chemistry, Chinese Academy of Sciences, P.O. Box 2711, Beijing 100080, China b Leibniz Institute of Polymer Research, Hohe Street 6, D-01069 Dresden, Germany c Centre for Advanced Materials and Technology, School of Aerospace, Mechanical and Mechatronic Engineering (J07), The University of Sydney, Sydney NSW 2006, Australia Received 5 March 2004; revised 18 November 2004; accepted 18 November 2004
Abstract The fracture surface morphology of short fiber reinforced thermoplastics (SFRTs) has often been used to assess qualitatively the degree of fiber–matrix interfacial adhesion. Mechanical properties such as tensile strength, fracture toughness and failure strain, etc. are then correlated with the morphology. Fracture surfaces showing fibers surrounded by a large amount of matrix material is commonly regarded as indication of strong fiber–matrix interfacial adhesion while smooth fibers are characteristic of weak interfacial adhesion. Many experimental results of SFRTs have been so interpreted. However, it is shown in this paper that strictly speaking, such interpretations are generally incorrect. Moreover, the amount of matrix material does not provide a quantitative measure of the adhesion. Correct implication of the morphology of fracture surfaces is clarified. Short glass fiber reinforced polyamide 6,6/polypropylene (PA 6,6/PP) blends toughened by rubber are employed as examples for SFRTs since the PA 6,6/PP blend system by changing PA 6,6 concentration in the matrix blend represents a wide range of matrix materials. It is demonstrated that the fracture surface morphology of such composites is dependent on both fiber–matrix interfacial adhesion strength and matrix shear yield strength. Consequently, tensile failure strain is well correlated with the post-mortem fracture surface morphology of these SFRTs. q 2004 Elsevier Ltd. All rights reserved. Keywords: A. Glass fibres; B. Interface/interphase; Thermoplastics
1. Introduction Short fiber reinforced thermoplastics (SFRTs) are very attractive because of their ease of fabrication, economy and superior mechanical properties [1–6]. Characterization of the morphology of the post-mortem fracture surfaces of SFRTs is often conducted with scanning electron microscopy (SEM). These morphologies provide direct evidence to study the strengthening and toughening mechanisms of these composites. It has been reported that fracture surfaces of SFRTs containing fibers surrounded by a substantial amount of matrix material imply strong * Corresponding author. Address: Technical Institute of Physics and Chemistry, Chinese Academy of Sciences, P.O. Box 2711, Beijing 100080, China. Tel.: C86 10 62659040; fax: C86 10 62564049. E-mail address: [email protected] (S.Y. Fu). 1359-835X/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.compositesa.2004.11.005
interfacial adhesion between fibers and matrix [7–9]. Conversely, smooth fiber surfaces signify weak interfacial adhesion between fibers and matrix [7–10]. The morphology of the fracture surfaces has often been used to correlate with the mechanical properties. Alsewailem concluded in his PhD project [7], from observation of the fracture surfaces with fibers surrounded by a large amount of matrix material, that addition of rubber to nylon 66 led to strong glass fiber–matrix interfacial adhesion, which in turn increased the impact strength. Nair et al. [8] attempted to link the fracture toughness of glass fiber reinforced rubbertoughened nylon 6,6 composites to their post-mortem fracture morphologies, such that clean pullout fibers on the fracture surfaces implied weak fiber–matrix interfaces while profound matrix-adhered fibers denoted strong fiber– matrix interfaces. It was then claimed that when the fiber– matrix interface was strong, the composites were tougher
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than the unreinforced polymer matrices; when the fiber– matrix interface was weak, the interface had little effect on the toughness. On the other hand, Stamhuis [10] compared the fracture surface morphologies of PP/GF/EPDM and PP/ GF/SBS composites and found no differences due to the two ealstomers (EPDM and SBS). However, better impact behavior was found for the PP/GF/EPDM composite. Further evidence is provided by Cho and Paul [11] who studied the mechanical properties and morphology of glass fiber reinforced ABS toughened nylon 6 composites. The ABS material used was a high rubber, emulsion-made product. Smooth fiber surfaces were observed for the composite without ABS. When the matrix was blended with ABS, there appeared to be some matrix material adhered to the fiber surfaces and the amount seemed to vary directly with rubber content. If the fracture surface morphology—interfacial adhesion—composite mechanical properties correlation is correct, the composite strength would increase with rubber content due to the enhanced interfacial adhesion. However, it was reported [11] that the composite strength actually decreased with increasing ABS content. Therefore, there is concrete evidence from previous studies to show the limitation of linking mechanical properties to post-mortem fracture surface morphology. The purpose of the present work is to clarify the implication of the morphology of the fracture surfaces of SFRTs. Tensile failure strain is then correlated with the post-mortem fracture surface morphology. Polyamide 6,6 (PA 6,6)/polypropylene (PP) blends were used as matrices since the PA 6,6/PP blend system by changing the PA 6,6 concentration in the matrix represented a wide range of matrix materials. A mixture of 10 wt% styrene–ethylene– butylene–styrene (SEBS) and 10 wt% maleic anhydride (MA) grafted SEBS (SEBS-g-MA), which was proven in earlier studies to be effective both as a compatabilizer and a toughener for PA 6,6/PP blends [12,13], was used to compatabilise the matrices. Short glass fibers (SGF) purchased from Owens Corning (123D) were used to reinforce the PA 6,6/PP blends.
2. Experimental details 2.1. Materials and processing Polyamide 6,6 (PA 6,6) resin (Vydyne 21) was obtained from Monsanto Chemical Company (USA) and polypropylene (PP) (Propathene GSE52) was supplied by ICI, Australia. The triblock copolymer SEBS (Kraton G1652) and its maleated version (Kraton F1901X) were bought from Shell Chemical. 123D short glass fibers of a length of 4 mm and an average diameter of 10 mm were purchased from Owens Corning. PA compatible sizing has been made by the producer on the surfaces of 123D glass fibers for reinforcing PA and their blends where hydrolysis resistance is not an end use requirement.
Polymer materials were first dry blended simultaneously at various weight ratios (100/0, 75/25, 50/50, 25/75 and 0/100) of PA 6,6/PP, 80 wt% PA 6,6/PP blend plus a mixture of 10 wt% non-maleated-SEBS and 10 wt% maleated SEBS copolymer. Then, short glass fibers were pre-compounded together with the dry blended polymer matrices at a weight ratio of 40/60 glass/polymer. Compounding of the materials was carried out using a high shear-rate twin-screw extruder (Werner and Pfleiderer ZSK 30) at a temperature range of 260–280 8C. Materials were injection molded into 3 mm thick tensile bars (ASTM D638 type I) by an injection molding machine (BOY 22S Dipronic). All materials were dried at 80 8C in an oven over night prior to compounding and injection molding. Samples from polymer materials were also prepared under the similar processing conditions so that the tensile data from these samples can be used as that of the matrix blends in the composite samples. 2.2. Tensile testing and post-mortem fracture morphology Tensile tests were conducted at room temperature at a crosshead rate of 5 mm/min using an Instron Model 5567 computer-controlled testing machine. Six samples for each composition were tested immediately after preparation of the samples. The samples were not humidity conditioned before testing. Scanning electron microscopy (SEM) examination was conducted on the fracture surfaces of the samples. Before SEM examination, samples were gold sputtered by a JEOL JFC-1100E surface coating machine and viewed with a JEOL JSM-820 SEM. 2.3. Fiber dispersion, mean fiber length and critical fiber length Optical micrographs of the fiber dispersion across the sections of the samples perpendicular to the flow direction were taken after the cross-section were polished. Fiber lengths were determined by burning off the polymer matrix in a furnace at 550 8C for 1 h. The fibers were then dispersed with water on glass slides so that their lengths could be determined microscopically after the water evaporated. About 600–1000 fibers were measured for each composition to get the mean fiber lengths. Critical fiber lengths were approximately obtained by doubling the average of the maximum pull-out fiber lengths from the fracture surfaces of the samples for each composition since the maximum fiber pull-out length should be equivalent to half of the critical length in principle. Pull-out lengths of over 30 longest pulled-out fibers (shown in Fig. 1) were measured and the 10 maximum fiber pull-out lengths were averaged for each composition as half of the critical fiber length. By the way, another method can also be employed to estimate the critical fiber length. First, all the pullout lengths of the fibers, which vary from 0 to half critical length, are
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Fig. 1. SEM micrographs of fiber pullout at the fracture surfaces of 40 wt% SGF reinforced PA 6,6/PP composites. (A) 100% PA 6,6, (B) 75% PA 6,6, (C) 50% PA 6,6, (D) 25% PA 6,6 and (E) 0% PA 6,6 in the matrix.
estimated, then the average of all the measurements is multiplied by 4 to get the critical fiber length. Similar results can be obtained.
3. Results and discussion 3.1. Effect of PA 6,6 concentration in the matrix blend on interface adhesion strength Fig. 2 shows the tensile strength of SGF reinforced PA 6,6/PP blends. It can be seen that the composite tensile strength increases with increasing the PA 6,6 concentration in the matrix system. When the PA 6,6 concentration is equal to or higher than 75 wt% in the matrix blend, the composite strength is much higher than for other PA 6,6 concentrations. Table 1 indicates that the mean fiber length decreases with increasing the PA 6,6 concentration in the matrix system. The increase in PA 6,6 concentration in
Fig. 2. Effect of the PA 6,6 concentration in the PA 6,6/PP matrix system on tensile strength of 40 wt% SGF reinforced PA 6,6/PP composites.
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Table 1 Various quantities versus PA 6,6 concentration in the matrix blend PA 6,6/matrix (PA 6,6CPP) (%)
Fiber volume fraction (%)a
Mean fiber length (mm)
Critical fiber length (mm)
Interfacial adhesion strength (MPa)
Matrix tensile yield strength (MPa)
Matrix shear yield strength (MPa)
100 75 50 25 0
0.220 0.213 0.205 0.198 0.190
182.3 209.0 268.6 324.4 332.2
260.2 275.8 430.1 650.0 950.6
38.43 36.26 23.25 15.38 10.52
40.05 32.10 20.93 14.98 19.69
20.03 16.05 10.47 7.49 9.85
a
The densities of PA 6,6, PP, SEBS and glass fibers used for the computation of fiber volume fraction are, respectively, 1.13, 0.91, 0.88 and 2.55 g/cm3.
the matrix blend corresponds to an increase in fiber volume fraction as shown in Table 1. It is well known that the viscosity of polymers can be significantly increased by the addition of fibers. Therefore, the increase of fiber volume fraction would lead to lower average fiber length due to higher viscous forces exerting on fibers. On the other hand, as fiber volume fraction increases, fiber–fiber interaction would also cause more damage to fiber length [14]. It was
shown previously [15,16] that the tensile strength of short fiber reinforced thermoplastic composites increased with both fiber–matrix interfacial strength and mean fiber length since fiber orientation and dispersion are approximately the same for the investigated cases (please see Fig. 3 in which fibers are uniformly dispersed and preferentially aligned along the flow direction, namely sample axis direction in all the investigated samples). It has been shown in Table 1 that
Fig. 3. Optical micrographs of the cross-sections of the composites. (A) 100% PA 6,6, (B) 75% PA 6,6, (C) 50% PA 6,6, (D) 25% PA 6,6 and (E) 0% PA 6,6 in the PA 6,6/PP matrix system.
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mean fiber length decreases with increasing the PA 6,6 concentration in the matrix system, this would lead to reduction in composite tensile strength. Nonetheless, it can be observed that the composite strength increases with the increase of the PA 6,6 concentration. It can thus be inferred that the fiber–matrix interfacial adhesion strength must increase with the PA 6,6 concentration in the matrix system. This is further confirmed below. It is shown that the critical fiber length decreases as the fiber–matrix interfacial adhesion strength increases [15,16]. The relationship between critical fiber length (lf) and interfacial adhesion strength (ti) is given by the wellknown equation below [17]: lc Z rf sfu =ti
(1)
where rf is fiber radius and sfu fiber fracture strength. The fracture strength of fibers used here is 2 GPa [18]. Some fractured fibers can be observed at the fracture surfaces of the composites (Fig. 4), indicating that longest fibers at
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the fracture surfaces should have approximately critical fiber lengths because longer fibers would be fractured. Thus, the approximation of the critical fiber lengths can be obtained from Fig. 1 for various matrix materials and then, the interfacial adhesion strength can be estimated according to Eq. (1). The results are given in Table 1. Indeed, it is noted that the critical fiber length increases with the decrease of the PA 6,6 concentration in the matrix system and hence, the interfacial adhesion strength increases with increasing the PA 6,6 concentration. 3.2. The post-mortem morphology of the pullout fibers at the fracture surfaces of SFRTs It has been shown above that the interface adhesion strength increases with the increase of the PA 6,6 concentration in the matrix system. As a result, it can be conjectured that the amount of the matrix material surrounding the fibers standing proud at the fracture surfaces
Fig. 4. SEM micrographs of the fracture surfaces of the composites. (A) 100% PA 6,6, (B) 75% PA 6,6, (C) 50% PA 6,6, (D) 25% PA 6,6 and (E) 0% PA 6,6 in the PA 6,6/PP matrix system.
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Fig. 5. The morphologies of the fracture surfaces of 40 wt% SGF reinforced PA 6,6/PP composites. (A) 100% PA 6,6, (B) 75% PA 6,6, (C) 50% PA 6,6, (D) 25% PA 6,6 and (E) 0% PA 6,6 in the PA 6,6/PP matrix system.
should also increase with the increase of the PA 6,6 concentration in the matrix system according to the common belief on the morphology of the pullout fibers. However, Fig. 5 shows that the medium PA 6,6 concentrations (75, 50 and 25%) in the matrix system correspond to greater amounts of the matrix materials surrounding the fibers rather than for the 100% PA 6,6 case. The micrographs in Fig. 5 are arbitrarily selected for the observation of the morphology of the fiber surfaces. Micrographs containing a number of fibers with the similar morphology of fiber surfaces are available in Fig. 4 with a lower magnification. Thus, Fig. 5 can be regarded to represent the observation of the fiber surfaces at the whole fracture surfaces. Consequently, the amount of the matrix material does not provide a quantitative measure of the adhesion and the fiber–matrix interface adhesion cannot be simply linked to the post-mortem morphology of the pullout fibers on the fracture surfaces. The traditional interpretations presented previously for the morphology of the fracture surfaces of SFRTs are thus incorrect.
Table 1 shows the values for the matrix tensile yield strength obtained from tensile testing of matrix materials. Based on Mohr’s theory, the matrix shear yield strength must be about 50% the matrix tensile yield strength and the results for the matrix shear yield strength are then given in Table 1. It can be seen that the interfacial adhesion strength is obviously higher than the matrix shear yield strength except for the case of the 0% PA 6,6 in the matrix system where the interface adhesion strength is close to the matrix shear yield strength. It can then be reasonably inferred that the matrix shear yielding and then failure would occur before the interface failure except for the case of the 0% PA 6,6 in the matrix system. Thus fiber surfaces adhered with matrix materials have been observed except for the 0% PA 6,6 in the matrix system as shown in Fig. 5. For the SFRT containing no PA 6,6, fiber surfaces are quite smooth because interface debonding would take place before occurrence of matrix shear yielding due to the fact that the interface stress at fiber ends must be quite higher than
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Fig. 7. Stress–strain curves for 123D SGF reinforced PA 6,6/PP composites. Fig. 6. Effect of the PA 6,6 concentration in the matrix system on tensile failure strain of 40 wt% SGF reinforced PA 6,6/PP composites.
the average interface shear stress [19,20] while the interface adhesion strength is very close to the matrix shear yield strength. Consequently, the post-mortem morphology of the fracture surfaces is determined by both interface adhesion strength and matrix shear yield strength. 3.3. Correlation of the post-mortem morphology of the pullout fibers with composite failure strain Comparison of Figs. 5 and 6 shows that the changing trend of the morphology is the same as that for the failure strain with the PA 6,6 concentration in the matrix system. At the one extreme of 0% PA 6,6 (namely 100% PP) where interface adhesion is weakest and fiber surfaces are smooth, the tensile failure strain is the lowest. At another extreme of 100% PA 6,6 where interface adhesion is strongest but there is some amount of matrix adhering to fiber surfaces, the tensile failure strain is higher than for the 0/100 PA 6,6/PP SFRP but lower than for other composites. While between the two extremes, substantial amounts of matrix material surrounding fiber surfaces are observed, the failure strains are higher than for the two extremes. From the tensile curves of the composites shown in Fig. 7, it can be seen that the tensile stress–strain curves show linear deformation at lower stresses and non-linear deformation at higher stresses as having been shown previously [4]. The deviation of the curves from linear to non-linear deformation corresponds to either interfacial debonding or matrix shear yielding deformation. For the PA 6,6 based composite (100% PA 6,6 in the matrix system), the tensile curve deviates from linear to non-linear behavior at a relatively higher stress than for other composites but the deformation from occurrence of non-linear deformation to composite fracture is relatively small (namely a small matrix shear yielding deformation) when compared with other composites, bringing about a relatively small failure strain and thus a small amount of matrix material adhering to fiber surfaces because of the small matrix shear yielding
deformation. For the SFRT containing no PA 6,6 with a low interfacial adhesion strength, interfacial debonding would occur at a relatively low stress and the non-linear deformation does not last long to composite fracture due to the weak interface adhesion, leading to the lowest failure strain of the composite. This is because interface debonding takes place before matrix shear yielding, no or little matrix material is adhered to fiber surfaces and thus smooth fiber surfaces are observed. For the SFRT containing 75, 50 and 25% PA 6,6 in the matrix system, the non-linear deformation is relatively large, resulting in relatively large failure strains (corresponding to relatively large matrix deformation and thus relatively large amounts of matrix material adhering to fiber surfaces). Hence, in summary, the tensile failure strain can be well linked to the post-mortem morphology of the pullout fibers on the fracture surfaces. 4. Conclusions The implications of the post-mortem morphology of the pullout fibers on the fracture surfaces of SGF reinforced PA 6,6/PP blends as an example of SFRTs are discussed. The results have shown that the fiber–matrix interface adhesion cannot be simply linked to the post-mortem morphology of the pullout fibers on the fracture surfaces. The traditional interpretations presented previously for the morphology of the fracture surfaces of SFRTs are thus incorrect. The observation of the fibers surrounded by a great amount of matrix material at the fracture surfaces does not certainly denote a strong fiber–matrix interfacial adhesion. Moreover, it is displayed that the post-mortem morphology of the pullout fibers on the fracture surfaces is determined by both interface adhesion strength and matrix shear yield strength. Finally, it is observed that the tensile failure strain can be well linked to the post-mortem morphology. Acknowledgments This work was funded by the Talent’s Program, Chinese Academy of Sciences and the Australian Research Council
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on ‘Short Glass Fiber Reinforced Toughened Polymer Blends’. YWM was supported by an Australian Federation Fellowship. SYF is grateful for the support of the Alexander von Humboldt Foundation and the Australian Research Council.
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