Macromolecules 2007, 40, 5953-5958
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In Situ Compatibilizer Reinforced Interface between an Amorphous Polymer (Polystyrene) and a Semicrystalline Polymer (Polyamide Nylon 6) Yongsok Seo,*,† Jiseok Lee,† Tae Jin Kang,† Hyoung Jin Choi,*,‡ and Jinyeol Kim§ Intellectual Textile System Research Center (ITRC) and School of Materials Science and Engineering, College of Engineering, Seoul National UniVersity, Shillim9dong 56-1, Kwanakgu, Seoul, Republic of Korea 151-744, Department of Polymer Science and Technology, Inha UniVersity, Yonghyun4dong, Namku, Inchon, Republic of Korea, 402-751, and School of AdVanced Materials Engineering, Kookmin UniVersity, Chongnungdong, Sungbukku, Seoul, Republic of Korea 133-791 ReceiVed February 10, 2007; ReVised Manuscript ReceiVed May 30, 2007
ABSTRACT: We present an investigation of the reinforcement of the interface between a flexible amorphous polymer (polystyrene, PS) and a semicrystalline polymer (a polyamide, Ny6). Poly(styrene-co-maleic anhydride) was used as the compatibilizer. Fracture toughness was measured using the asymmetric double cantilever beam test (ADCB). For bonding temperatures above 190 °C, the adhesion strength was found to increase with bonding time, pass through a peak value, and then reach a plateau. The fracture toughness increased with increasing bonding temperature, passed through a peak near 200 °C, and then decreased with further increase of the bonding temperature. This behavior was more obvious for an amorphous polymer/semicrystalline polymer pair than for a pair of semicrystalline polymers. The variation of the fracture toughness with bonding time and temperature can be plausibly explained in terms of two different failure mechanisms: adhesive failure at the interface for short bonding times and when the bonding temperature is low, and for longer bonding times and at high temperatures, cohesive failure between chains at the interface and the bulk PS due to decreased chain entanglement. As long as the compatibilizer molecular weight is sufficiently large that entanglements form between the matrix polymers, common features of the fracture mechanisms depending on the bonding conditions can be outlined that are independent of the crystallinity of the polymer.
Introduction The interfacial properties and the properties of the component polymers are the strongest influences on the final physical and chemical properties of polymer blends. Since most polymer pairs are thermodynamically immiscible, their interfacial adhesion is quite poor due to poor mutual diffusion of molecules at the interface.1 To fabricate a useful blend, the addition of a compatibilizer at the interface is normally required.1,2 The addition of a compatibilizer such as a block copolymer is only useful if the compatibilizer is available. If no compatibilizer is available, in situ reactive compatibilization is another approach that can be used to improve the interfacial properties. Compatibilizer produced in situ at the interface can transfer the stress larger than the crazing stress of one of two matrix polymers. A compatibilized polymer blend with high adhesion strength can then be achieved because plastic deformation occurs in the more ductile polymer and the asymmetric craze propagates ahead of the crack.3,4 The mechanical role of a compatibilizing copolymer at the interface between two glassy polymers is now relatively well understood. As long as the copolymer’s molecular weight is sufficiently high that entanglements of the two matrix polymers form on both sides of the interface and the number density of the copolymer is sufficiently high, the presence of the com* To whom correspondence should be addressed. E-mail: (Y.S.)
[email protected]; (H.J.C.)
[email protected]. † Intellectual Textile System Research Center (ITRC) and School of Materials Science and Engineering, College of Engineering, Seoul National University. ‡ Department of Polymer Science and Technology, Inha University. § School of Advanced Materials Engineering, Kookmin University.
patibilizer at the interface provides high interfacial adhesion strength. If the areal density of the copolymer at the interface is insufficient to sustain a crazing stress, fracture occurs at the interface as a result of chain scission.5-8 Few studies of semicrystalline polymer interfaces have been carried out until very recently. The study of semicrystalline polymers is more complicated because they are typically two-phase materials, so their deformation properties depend on the crystallization kinetics and the resulting nonequilibrium properties (cocrystallization).9-12 The most heavily investigated system is the interface between polypropylene (PP) and polyamide 6 (Ny6) reinforced by reactive blending.11,12 The basic factors controlling the mechanical properties of the interface are the same as those for glassy polymers, but the additional effects of cocrystallization and microstructural effects on semicrystalline polymer pairs should be taken into consideration.12-14 Thus, the measured fracture toughness is expected to vary significantly with the bonding conditions (temperature and time). We have previously investigated PP/Ny6 interfaces reinforced with in situ copolymer formation.12 Maleic anhydride grafted polypropylene (MAPP) was premixed with the polypropylene and then adhered to the Ny6 surface. The fracture toughness was found to increase with the bonding time, pass through a peak, and then reach a plateau. The presence of a different crystalline phase (β-phase) in the PP sample was suspected to be the reason for the large increase in fracture toughness, but this could not be confirmed with wideangle X-ray spectroscopy.6,12,13 The dependence of the fracture toughness on the bonding time can be explained in terms of two fracture mechanisms, adhesive failure, and cohesive failure. However, the high fracture toughness for a particular temperature remains to be explained.
10.1021/ma070354b CCC: $37.00 © 2007 American Chemical Society Published on Web 07/17/2007
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In a separate study, we verified experimentally that the appearance of a maximum in the fracture toughness is not because of variation of the amount of compatibilizer formed at the interface nor because of functionalized molecular diffusion to the interface.15 To obtain these results, we used a surface functionalization method using ion-beam irradiation in an oxygen environment, which produces some reactive functional groups on the PP surface at a relatively shallow depth. The fracture toughness was found to exhibit almost identical behavior: the fracture toughness passes through maxima at 200 °C and at a bonding time of 60 min. The difference between this temperature for maximum toughness and that for the maleic anhydride grafted polypropylene (MAPP + PP)/Ny6 system was ascribed to the limited availability of functionalized molecules from the bulk.15 Thus, we concluded that the results for the surface-functionalized PP/Ny6 system are reaction-rate limited whereas for the (MAPP + PP)/Ny6 system they are favorably limited by diffusion.16 From the experimental results reported so far, we conclude that as long as the compatibilizer molecular weight is sufficiently large that entanglements form between the matrix polymers, common features of the fracture mechanisms depending on the bonding conditions can be outlined that are independent of the crystallinity of the polymer. To verify this conclusion, we have carried out a further investigation of the fracture toughness of the interface between an amorphous polymer (polystyrene) and a semicrystalline polymer (Ny6). Poly(styrene-co-maleic anhydride) with a high molecular weight was used in the in situ reactive compatibilization. A fracture mechanism that is universally applicable to polymer interfaces was sought for the in situ reactive compatibilization process. The effect of bonding conditions on the fracture energy was also studied by varying the bonding temperature and bonding time.
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Figure 1. Schematic diagram of the ADCB test.
after an hour when there was no further increase in the crack length. Several images were taken on the same sample and the same procedure was applied to different samples to get the reproducible data. Boucher et al.8 reported that the ADCB test yielded reliable values of the energy of adhesion, Gc, if two precautions were taken. First, the samples had to be asymmetric because the different mechanical properties of the two polymers might induce various modes of fracture. They also noted that varying the ratio of thickness changed the amount of the KII mode in the fracture process19,20 and that if the fracture tended to deviate into the more ductile material, the measured energy release rate could increase significantly, leading to substantial errors in the evaluation of Gc.17 To minimize contributions of the second component, all of our samples were made with a thickness ratio hPS/htot of 0.67 because Gc had a minimum value at a ratio between 0.55 and 0.7.21 In this system, Young’s moduli of PS and Ny6 are 1.5 and 2.05 GPa, respectively. Since the crack length ahead of the blade, a, was less than 10hPS for most of our samples, the following equation derived by Boucher et al.8 based on calculations by Kanninen,20 whose assumption was that the finite elasticity of the material ahead of the crack tip required correction factors for small crack lengths, was used: Gc )
Experimental Section Materials. Materials employed in this study were commercial polyamide (Ny6) and a polystyrene (PS). Polystyrene was supplied by Kumho Petrochemicals (Korea). The weight-average molar mass was 2.8 × 105 g/mol and the polydispersity index was 2.4. Ny6 was a Kolon product (KN171), whose weight-average molar mass was 8.5 × 104g/mol and the polydispersity index was 3.5. Commercially available poly(styrene-co-maleic anhydride) (PSMA), added as a compatibilizer, was purchased from Aldrich in pellet form. It contains 7 wt % maleic anhydride group. The weightaverage molar mass was 2.34 × 105 g/mol and the polydispersity index was 2.3. Pellets of all polymers were dried for 24 h in a vacuum oven at 100 (Ny6) and 80 °C (PS and PSMA). Blending of PS and PSMA (97:3 weight ratio) was done in an internal mixer at 200 °C. Samples were made by compression molding at 160 and 240 °C for PS blend and Ny6, respectively. The PS blend strips (2 cm × 4 cm) were clamped with Ny6 strips (2 cm × 4 cm) in an airtight molder under slight pressure. The mold was heated in a temperature-controlled furnace between 180 and 220 °C. We did not use bonding temperatures higher than that of the melting temperature of Ny6 (225 °C) because the interface becomes unstable when both samples are in the melt state. The mold was slowly cooled to room temperature in air. All the samples were stored in a desicator for 24 h prior to fracture test. Measurement of the Fracture Toughness. The fracture toughness was measured using the asymmetric double cantilever beam (ADCB) test because it has been shown to be a reliable test for the fracture toughness of a polymer interface.15-19 Details of this test are shown in Figure 1; a blade of thickness ∆ was inserted at the interface between PS (+PSMA) and Ny6 and was pushed into the sample. Because the PS side at the experimental thickness was transparent, a video camera was used to measure the crack length exactly. An image of the region ahead of the blade was recorded
EPShPS3ENy6hNy63 3∆2 8a4 EPShPS3RNy62 + ENy6hNy63RPS2
(1)
where Ei and hi denote the Young’s modulus and the thickness of material i, respectively, and ∆ is the thickness of the blade. Ri is the correction factor for material i and is given by
(
Ri ) 1 + 1.92
()
hi hi + 1.22 a a
2
+ 0.39
( ))( hi a
3
/ 1 + 0.64
)
hi a
(2)
Surface Characterization. Scanning electron microscopy (SEM) observations of the samples were performed on a Hitachi S-2500C. Fractured surfaces were coated with gold in an SPI sputter coater. The morphology was determined using an accelerating voltage of 15 keV. Chemical components on the fractured surfaces were analyzed by XPS. XPS spectrum was recorded by Surface Science 2803-S spectrometer (hν ) 1.5 keV). The basic pressure of 2 × 10-10 Torr was maintained during the analysis. Energy resolution of 0.48 eV was kept. The XPS spectra were referenced to the main component of the C 1s peak of PSMA at 284.6 eV of binding energy, O 1s peak at 532 eV for CdO peak and 533.6 eV for C-O peak and N 1s peak of Ny6 at 399.7 eV. The overlapping peaks were resolved by the peak synthesis method based on Gaussian peaks.
Results and Discussion The Variation of Fracture Toughness with Bonding Time and Temperature. The chemical reactions occurring at the interfaces of reactively blended polymers are controlled either by the rate of diffusion or by the rate of reaction.15 The parameters affecting the rates of diffusion and reaction are the temperature and the initial concentration of the functional groups on the polymer surface. Since we cannot precisely control the concentration of functional groups at the polymer interface
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Figure 2. Variation of the fracture toughness of PS(+PSMA)/Ny6 interface with bonding time. Bonding temperatures were (b) 180, (O) 190, (1) 200, (4) 210, and (9) 220 °C, respectively.
Figure 3. Variation of the fracture toughness between PS(+PSMA) and Ny6 with bonding temperatures. Bonding times were (b) 30, (O) 60, (1) 75, (4) 90, and (9) 120 min. The lines are guides for the eyes.
because of the premixing of SMA with PS, we examined the effects of varying the bonding temperature and those of varying the bonding time at constant bonding temperature, both of which definitely affect the diffusion of molecules and the reaction at the interface. We emphasize again that the bonding temperature never goes over 225 °C of Ny6 melting temperature since the interface position is not stationary at a fixed position if both phases are in the melt state. Hence the Ny6 side simply acts as a hard wall with immobile functional groups. Figure 2 shows the variation of the fracture toughness with bonding time. As was found for the adhesion between two semicrystalline polymers (polypropylene and nylon-6, PP/Ny6),12 several facts are worthy of note. First, the interfacial fracture toughness for each temperature increases with the bonding time. Up to a bonding time of 30 min, the fracture toughness increases very slowly, which indicates that not much reaction occurs at the interface, so some induction annealing time is required.15 For longer bonding times, the fracture toughness passes through a maximum and/or reaches a plateau value depending on the bonding temperature. Second, the fracture toughness at the same bonding time increases with bonding temperature, reaching its highest value at 200 °C, and then decreases with further increases in the temperature. Third, there is a clear maximum in the fracture toughness at a bonding time of around 60 min for all bonding temperatures. Fourth, the adhesion strength at the interface increases more rapidly at higher bonding temperatures. Since the cooling conditions for all of the samples were the same, the early increase in the interfacial adhesion with bonding time is mainly due to an increase in the number of intermolecular reactions at the interface. These results are very similar to those for reactive interfacial adhesion between semicrystalline polymers (PP/Ny6),12,15 and a semicrystalline polymer (MAPP) and a thermotropic liquid crystalline polymer.16 As we have pointed out previously, the areal density of the produced copolymers cannot increase indefinitely because diffusion of functional molecules from the bulk to the interface becomes more and more difficult as reacted molecules occupy the interfacial area.12 As the intermolecular reactions between the PSMA and Ny6 molecules proceed, the interface becomes overcrowded with the graft copolymers generated at the interface. The fracture toughness increases with the areal density of the graft copolymers. We also reported that the fracture toughness at a PP/Ny6 interface compatibilized with maleic anhydride grafted polypropylene (MAPP) increases with bonding time until the critical areal density of the graft copolymers reaches a saturation point, at which further reactions are hampered near the interface by the steric hindrance produced
by the high areal densities of the copolymers.12 Though further diffusion of PSMA molecules from the deep bulk side to the interface is difficult due to the hindrance, further reaction between the functional groups of PSMA and Ny6 can happen because the PSMA polymer chains are multifunctional. Hence the type of grafting depends on temperatures and time. As Char and Lee22 explained, the situation may evolve from one graft per SMA chain to several grafts per chain reducing hence greatly the molecular weight of the SMA chain portion (loop or chain end) able to entangle with PS homopolymer. As a result, the molecular weight of the “effective brush” at the interface decreases.22 Microphase separation between the molecules in the bulk and those at the interface would then intervene above a certain saturation threshold.23 This point will be explained again later. The variation of the fracture toughness with bonding temperature is shown in Figure 3. The growth rate increases with temperature up to 200 °C. There is a clear maximum in the fracture toughness near 200 °C, which decreases at higher bonding temperatures. Morphology and Fractured Surfaces. As long as the copolymers (or the block copolymers added as a compatibilizer) at the interface have sufficiently high molecular weights to become fully entangled with the matrix polymers on both sides of the interface, fracture occurs by chain scission when the areal density of the copolymers at the interface is low.9 If the areal density of a long compatibilizer is above some critical value, the adhesive strength is high enough to withstand the fracture stress. The fracture is then determined by the cohesive strength of the polymers on either side, depending on the states of the individual chains. Thus, the failure at the interface proceeds as a result of adhesive failure or cohesive failure, i.e., failure of the interface proceeds by the mechanism with the lower intrinsic failure energy. In the present study, cohesive failure proceeds only in the PS side because the crazing stress of Ny6 is higher than that of PS. As already mentioned, all the deformation is, thus, in the PS side and the NY6 side acts like a solid wall with functional groups. The locus of the failure was investigated microscopically. Figure 4 shows SEM micrographs of the fractured surfaces on the PS side of the interface for a series of specimens bonded at a bonding temperature of 200 °C. For the short bonding time specimens, not enough reaction occurred at the interface. Hence, the fracture surface consists of large flat areas due to weak adhesion. As the bonding time increases, so does the adhesion strength, and a large fracture occurred for an bonding time of 90 min. This result shows that for a longer bonding time there is more crack propagation on the PS side of the interface and
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Figure 4. SEM photographs of the cleaved PS surfaces at a bonding temperature of 200 °C for different bonding times (×20000): (a) 30, (b) 60, (c) 90, and (d) 120 min.
Figure 5. Schematic representation of the variation of the interfacial reactions with the reaction time:15,22 (a) early period; (b) later period. Black lines are PSMA chains. Light gray lines are PS chains and dark gray lines in the lower side are Ny6 chains. Empty and filled circles represent the functional groups on the PSMA molecules (succinic anhydride groups) and Ny6 molecules (amine groups). Because of entanglement with other Ny6 and PS molecules on both sides of the interface, the adhesion strength increases with bonding time. The type of grafting depends on temperature and time. For longer bonding times, more graft copolymers are formed at the interface. As the reaction time increases, since the PSMA polymer chains are multifunctional, and that the situation may evolve from one graft per PSMA chain to several grafts per chain hence greatly reducing the molecular weight of the PSMA chain portion (loop or chain end) able to entangle with the PS homopolymer. As a result the molecular weight of the “effective brush” at the interface decreases. The cohesive strength then decreases with increases in the bonding time, and this occurs more rapidly at a higher bonding temperature.
a higher fracture toughness comes out. Since the cracks propagate in the material with a lower yield stress (i.e., in the polystyrene side) when the interface is strong, the fractured surface of the PS side becomes rougher and more rugged with the production of more copolymers at the interface as a result of the increased number of chemical reactions occurring for longer bonding times. As mentioned above, the PS and Ny6 used in this study had molecular weights much higher than the critical value for chain entanglement, so complete cohesive failure was observed in this case because of strong adhesion at the interface. Fracture Mechanisms at the Interface. In the PS/Ny6 system with added SMA, the amine end groups of Ny6 and the maleic anhydride groups of SMA easily react to form covalent bonds during processing.12 Interface is occupied by graft copolymers with grafting branches that are entangled with Ny6 molecules in the bulk and participate in the cocrystallization of Ny6 molecules. When a critical surface coverage of bonded molecules is attained, further diffusion of functionalized groups from the bulk to the interface becomes quite difficult.22 However, still more reaction can happen. Figure 5 shows a schematic representation of the variation of the interfacial reactions with the reaction time.22 Time dependence of the fracture toughness can be explained as follows. The initial increase in fracture toughness with the bonding time is obviously due to an increased number of reactions.23 Because of entangle-
ment with other Ny6 and PS molecules on both sides of the interface, the adhesion strength increases with bonding time. For longer bonding times, more graft copolymers are formed at the interface. Even though further diffusion of PSMA molecules from the deep bulk side to the interface is difficult due to the hindrance of preoccupied grafted molecules at the interface, more reaction between the functional groups of PSMA and Ny6 can happen because the PSMA polymer chains are multifunctional. We mentioned earlier that as the reaction time increases, the situation evolves from one graft per SMA chain to several grafts per chain reducing hence greatly the molecular weight of the SMA chain portion (loop or chain end) able to entangle with PS homopolymer and resulting in less entanglement of that PSMA molecule (thick black line in Figure 5) at the interface with other PS molecules. As a result, the molecular weight of the “effective brush” at the interface decreases. The cohesive strength then decreases with increase in the bonding time, and this occurs more rapidly at a higher bonding temperature. The overall fracture toughness increases initially with bonding time due to the resulting increased adhesive strength. For longer reaction times, the adhesive strength increases above the value of the cohesive strength, which decreases with reaction time. Hence there is a maximum in the variation of the fracture toughness with time. The temperature dependence can be explained similarly. At the low temperature of 180 °C, the reaction proceeds slowly. The concentration of grafted copolymers produced at the interface is not sufficient to produce a strong interface. Even though molecules with functional groups for reaction at the interface are found to have a strong tendency to move to the interface, the entropic constraints on the diffusion of functionalized molecules to the interface from the bulk are too large to be overcome by a specific energetic driving force between functional groups only. At the high bonding temperature, however, more enthalpic energy is provided for PSMA polymer molecules to overcome the entropic constraints. The situation then changes from one graft per PSMA chains to several grafts per chain reducing the molecular weight of the SMA chain portion able to entangle with the PS homopolymer. As a result, the molecular weight of the effective brush at the interface decreases. The entanglement of a PSMA molecule with the PS molecules decreases with temperature in spite of many copolymers at the interface, and so the fracture toughness becomes low. This outcome is summarized in Figure 6. The fracture toughness at first increased with annealing time because of the increased adhesive strength due to interfacial reactions. When the adhesive strength is lower than cohesive strength, the failure occurs first through the adhesive failure. If the cohesive strength becomes lower than the adhesive strength, failure occurs due to failure of cohesion between the molecules at the interface and PS molecules below the interface. If the fracture strength is plotted vs bonding time, whether a maximum appears or not depends on the relative magnitudes of the adhesive strength and
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Figure 7. N 1s XPS spectra of the cleaved surfaces annealed at 200 °C: (a) PS side; (b) Ny6 side. Bonding times were 30 min (- -) and 90 min (-).
Figure 6. Schematic representation of the locus of failure of PS(+PSMA)/Ny6 interface. When the adhesive strength (-) is lower than cohesive strength (- -), the failure occurs first through the adhesive failure. But after the former becomes larger than the latter, the failure occurs through the cohesive failure. Bold lines implicate the overall path of the failure with time at different temperatures. The arrow indicates the fracture toughness variation with the bonding temperature after a long time (120 min). Table 1. Elemental Compositions of Fractured PS and Ny6 Surfaces Measured by XPS annealing time (at a constant bonding temp of 200 °C), min
PS %O
%N
%O
30 60 90 120
9.76 6.96 4.25 5.33
9.37 5.49 4.62 5.92
12.88 9.57 8.66 10.17
annealing temp (at a constant bonding time of 90 min), °C
PS %O
%N
%O
180 190 200 220
9.54 8.11 4.25 6.21
9.45 8.49 4.62 6.64
12.52 9.59 8.66 8.87
Ny6
Ny6
the cohesive strength. After sufficient time, both the adhesive strength and cohesive strength reach steady-state values. The temperature dependence shows a similar behavior. As the bonding temperature increases, more reactions occur faster. The adhesive strength increases more rapidly with the reaction time at high temperatures, and the cohesive strength also decreases more rapidly. Depending on the relative values of the adhesive and cohesive strengths, the total adhesion strength varies with the temperature (see the arrow in Figure 6). The elemental compositions of the fractured PS and Ny6 sides of the interface, as measured with XPS, support this explanation. The XPS results of the cleaved surfaces are listed in Table 1. On the PS side, the disappearance of oxygen functionalities (in the MA segments) is due to the transfer of PSMA chains. The crack propagates into the PS phase because of very strong adhesion at the interface, thus leading to the disappearance of oxygen moieties. In Table 1, the key elemental difference between PS (or PSMA) and Ny6 is the presence of nitrogen in Ny6 (Figure 7). In all of the tested fractured joints, we found no nitrogen on the PS side (Figure 7). The amount of nitrogen on the Ny6 side decreases at first with increases in the bonding time, implying the coverage of PS molecules on this side, and increases for longer bonding times. The amount of oxygen does
not change significantly after 120 min (not listed in Table 1). The initial decrease is obviously due to an increase in the number of PS molecules transferred to the Ny6 side with increases in the bonding time. The bonding time dependence for longer bonding times is related to the reaction of PSMA chains, as we explained before, decreasing molecular weight for the entanglements with other PS chains. Table 1 also lists the XPS results for the samples bonded for 90 min at various temperatures. The overall behavior is similar to the variation of the fracture toughness with bonding time. Conclusions In this study, we attempted to determine plausible fracture mechanisms at the polymer-polymer interface by investigating experimentally the effect of in situ reactive compatibilization on the fracture toughness of the interface between an amorphous polymer (PS) and a semicrystalline polymer (Ny6) compatibilized with PSMA. In general, the behavior of this interface was found to be similar to that of the interface between semicrystalline polymers (PP/Ny6) compatibilized by maleic anhydride grafted polypropylene (MAPP). The fracture toughness was found to increase with bonding time, pass through a maximum value, and then reach a plateau for bonding temperatures higher than 190 °C. The fracture toughness also increases with the bonding temperature, with a maximum near 200 °C and then decreases at higher bonding temperatures. Since the PSMA polymer chains are multifunctional, the type of grafting depends on temperature and time and the situation evolves from one graft per PSMA chain to several grafts per chains reducing greatly the molecular weight of the PSMA chain portion (loop or chain end) able to entangle with the PS homopolymer. As a result the molecular weight of the “effective brush” at the interface decreases. Increasing grafting density increases the adhesive strength whereas decreasing brush molecular weight decreases the cohesive strength. Therefore, the variation of the fracture toughness with bonding time and temperature can be plausibly explained in terms of two different failure mechanisms, i.e., adhesive failure at the interface for short bonding times and/or low bonding temperature and the failure of cohesion between the chains at the interface and the bulk of the lower modulus polymer (PS) for long bonding times and/or hightemperature due to decreased chain entanglements. When the cohesive strength becomes lower than the adhesive strength, fracture occurs by the cohesive failure in the PS side. Depending on the relative values of the cohesive and adhesive strengths, the total fracture toughness is decided. Therefore, there is an optimum bonding times and temperature for the fracture toughness which appears as a maximum in the variation of the fracture toughness with bonding time and temperature. This behavior is more obvious for an amorphous polymer/semicrystalline polymer pair than for semicrystalline polymer pairs
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because of the absence of cocrystallization in the amorphous polymer phase.12 Acknowledgment. This work was supported by KRF (041720060134 to Y.S.), the SRC/ERC program of MOST/KOSEF (R11 - 2005-065 to T.J.K.), KOSEF (R01-2006-000-10062-0 to H.J.C.), and MOCIE (N.13-5060) to J.K. References and Notes (1) Paul, D. R., Newman, S., Eds. Polymer Blends; Academic Press: New York, 1978; Vol. 1. (2) Paul, D. R., Bucknall, C. B., Eds. Polymer Blends; John Wiley and Sons: New York, 2000; Vols. 1 and 2. (3) Wool, R. P. Polymer Interfaces: Structure and Strength; Carl Hanser Verlag: New York, 1995. (4) Creton, C. In Polymer Surfaces and Interfaces; Richards, R. W., Peace, S. K.; John Wiley and Sons: Chichester, U.K., 1999. (5) Creton, C.; Kramer, E. J.; Brown, H. R.; Hui, C. Y. AdV. Polym. Sci. 2002, 156, 53. (6) Plummer, C. J. G.; Kausch, H. H.; Creton, C.; Kalb, F.; Le´ger, L. Macromolecules 1998, 31, 6164. (7) Norton, L. J.; Smigolova, V.; Pralle, M. U.; Hubenko, A.; Dai, K. H.; Kramer, E. J.; Hahn, S.; Berglund, C.; DeKoven, B. Macromolecules 1995, 28, 1999. (8) Boucher, E.; Folkers, J. P.; Creton, C.; Hervert, H.; Leger, L. Macromolecules 1997, 30, 2102. (9) Lin, J. J.; Silas, J. A.; Bermudez, H.; Miltam, V. T.; Bates, F. S.; Hammer, D. A. Langmuir 2004, 20, 5493.
Macromolecules, Vol. 40, No. 16, 2007 (10) Kim, H. J.; Lee, K.; Seo, Y.; Kwak, S.; Koh, S. Macromolecules 2001, 34, 2546. (11) Laurens, C.; Creton, C.; Leger, L. Macromolecules 2004, 37, 6814. (12) Seo, Y.; Tran, H. N. Polymer 2004, 45, 8573. (13) Kalb, F.; Leger, L.; Creton, C.; Plummer, J. G.; Marcus, P.; Malgalhaes, A. Macromolecules 2001, 34, 2702. (14) Benkoski, J. J.; Flores, P.; Kramer, E. J. Macromolecules 2003, 36, 3289. (15) Kim, H. J.; Lee, K.; Seo, Y. Macromolecules 2002, 35, 1267. (16) Seo, Y.; Ninh, T. H.; Hong, S. M.; Kim, S.; Kang, T. J.; Kim, H.; Kim, J. Langmuir 2006, 22, 3062. (17) Creton, C. F.; Kramer, E. J.; Hui, C. Y.; Brown, H. R. Macromolecules 1992, 25, 3075. (18) Brown, H. R. Macromolecules 1991, 24, 2752. (19) Brown, H. R. J. Mater. Sci. 1990, 25, 2791. (20) Kanninen, M. F. Int. J. Fract. 1973, 9, 83. (21) Lee, J. S. M.S. Thesis, Seoul National University, Seoul, Korea, 2006. (22) Lee, Y.; Char, K. Macromolecules 1994, 27, 2605. (23) Though the homopolymer chains are thus eventually expelled from the copolymer brush at the interface due to the high entropic energy barrier there, we doubt if normal PS molecules near or at the interface prefer to escape from the high energetic interface by their own diffusion. This needs to be studied further in the future. (24) Using the method by Boucher et al,8 the areal density of PSMA can be decided if the molar mass distribution of PSMA is narrow. The polydispersity index of PSMA is, however, not small, and reliable data were hard to obtained for our samples.
MA070354B
