Preparation of a structured acrylic impact modifier and its application in toughening polyamide 6.
It is well known that polyamide 6, is a notch-sensitive plastic. It is sensitive to crack propagation and becomes embrittled at low temperature or in dry ambient. Therefore, its impact resistance is poor and it tends to fail in a brittle manner. Many attempts have been made to improve the impact strength of PA 6 by adding low modulus toughening agents, i.e., rubbers, to decrease nolch-sensitivitv and increase toughness.
There are extensive technical literatures on rubber-toughened PA 6. Various types of rubber have been used, including styrene-elhylene/butylene-styrene block copolymer (SEBS) and/or SEBS grafted with maleic anhydride (SEBS-g-MA) (1-3) (EPR), and/or (EPR-g-MA) (4-13), styrene-acrylic acid copolymer (3), acrylonitrile-butadi-ene-styrene copolymer (ABS) (14), (15), polyethylene-octene copolymer (EOR) and/or EOR grafted with maleic anhydride (EOR-MA) (16), (17), epoxidized ethylene propylene diene rubber (18), ethylene propylene diene rubber grafted with maleic anhydride (EPDM-g-MA) (19), (20), EPDM grafted with styrene acrylonitrile copolymer (21), core-shell impact modifier (22), (23), polyvinyl acetate blended with ethylene-acrylic acid copolymer (24), ethyl-ene-acrylic acid copolymer (25), polybutadiene (26), natural rubber with maleic anhydride (NR-g-MA) (27), car-boxy lated styrene-butadiene rubber (28), acrylonitrile-bu-tadiene copolymer (29), carboxylated nitrile rubber (30), ultra-fine fully-vulcanized acrylate powdered rubber (31), and carboxylic styrene-butadiene ultra-fine full-vulcanized powdered rubber (32). Effective toughening agents must be functionalized or polar rubbers because PA 6 is a polar material. Therefore, it is plausible to increase interfacial adhesion or to obtain reactive blending between PA 6 and functionalized/polar rubber. Maleated rubber has been selected for this application. The grafted maleic anhydride reacted with the end group of PA 6 yielding an in situ graft copolymes on the interface between the two phases. This graft copolymer acted as an emulsifier, leading to a decrease in interfacial tension, particle-particle coalescence rate and rubber particle size. As a result, improvement in toughness was achieved to some extent. However, in the traditional blending methods size and dispersion of the rubber phase are difficult to control. Moreover, a disadvantage of rubber toughening is the significant reduction of modulus or stiffness that accompanies the addition of the soft rubber phase to the rigid polyamide. Materials with a better balance between stiffness and toughness are needed to extend their applications in various fields. Therefore, a kind of three layers core-shell particle structure with rubber phase core, rigid phase and functional monomer in shell could be designed. The rubber phase increases the impact strength; meanwhile, the rigid polymer phase offers a route to increasing the modulus. Functional monomer in shell may be chemically bonded to the matrix by reacting its functional groups with the amine groups of the PA 6. The rubber phase is surrounded by rigid polymer, thus, the core-shell modiiiers can flow unrestrictedly in PA 6 matrix. Poly(n-butyI acry!ate)/pol-y(methyl methacrylate-ro-acrylic acid), i.e., PBMA core-shell structured model with lightly cross-linked poly (BA) rubber phase as a core, rigid PMMA phase as a shell and acrylic acid in shell which is chemically bonded to the matrix by reaction its carhoxyl units with the amine groups of the PA 6 is selected. This article focuses on the preparation of acrylic core-shell particles prepared in a semicontinuous emulsion polymerization process and the investigation of the toughening effectiveness of these reactive core-shell modifiers on PA 6. An improved tough-mmstiffmm balance was achieved in PA 6 matrix by adding PBMA core-shell structured copolymer.
The initiator potassium persuifate (KPS) was obtained from Tianjin Chemistry Agent (Tianjin City, China). The anionic surfactant used in this study was Aerosol MA Series from Cytec, (Hevens City, The Netherlands). Tert-dodecylmercaptan (TDM; Fluka, Gillingham, UK) was used as supplied. All these materials were used without further purification. w-Bulyl acrylate (BA), acrylic acid (AA), and methyl methacrylaie (MMA) were purchased from Beijing Dongfang Chemical of China (Beijing City, China). Ally! methacrylate (ALMA), purchased from Tianjiao Chemical of China (Tianjin City, China), was used as received. The BA monomer was freed of inhibitor by washing with a 2 wt% NaOH solution; it was then washed with deionized water until the washed waters were neutral and finally dried with CaCI2 overnight, after which it was distilled under reduced pressure. MMA and AA monomers were purified by distillation under reduced pressure before use. Hydroquinone was used as an inhibitor of the latexes taken from the emulsion polymerization procedure. Deionized water was used throughout the study. PA 6 resin (1013B) was provided by Ube Kosan; Japan. All components were dried for at least 12 h at 80' C in a vacuum oven to ensure removal of adsorbed water.
Semicontinuous Emulsion Polymerization Process
Core-shell structured latexes were synthesized as 50% solid latexes via a two-stage semicontinuous emulsion polymerization. The weighed surfactant (0.5 g) and water (700 g) were added to a 3-L flanged reaction flask. The flow of nitrogen was started, and the water hatch temperature was kept at 78[degrees]C. During the following 30 min, the seed-stage BA monomer (50.0 g, 5 wt% of total monomer) was added to the surfactant solution and stirred for 10 min before KPS (2.15 g) dissolved in water (J00 g), which was added to initiate the reaction. The seed stage was 60 min. The growth stage involved three layers of preemulsified monomers, which were the first layer of preemulsified monomer of BA (750 g) with surfactant (9.4 g), ALMA (4.5 g) and weighed TDM (0.285 g), the second layer of preemulsified monomer of MM A (40 g) with surfactant (0.5 g) and the third layer of preemulsified monomer mixture of MMA (152.95 g) and AA (8.05 g) with surfactant (2 g), respectively. On the basis of the monomer weight three layers of preemulsified monomers were dropped into the flask successively at a constant rate over 3 h. KPS (0.215 g) dissolved in distilled water (50.0 g) was added to the reaction flask at 115, 175, and 235 min. After the completion of the addition of the growth-stage reactant mixture, a further 60 min was allowed before the latex was cooled to room temperature and filtered through a 53-/mi sieve to obtain the coagulate content. On the basis of the above latex preparation, contents of crosslinking agent and acrylic acid were regulated, respectively, while the seed stage was the same. To the variation of crosslinking agent content (ALMA), the growth stage involved three layers of preemulsified monomers, which were the first layer of preemulsified monomer of BA (750 g) with surfactant (9.4 g), ALMA (3.75 g, 5.25 g, 6 g, and 6.75 g, respectively, corresponding to 0.5, 0.7, 0.8, and 0.9% based on the first layer monomer's weight percent) and weighed TDM (0.285 g), and the contents of the second and the third layers were as same as above. To (he variation of the functional monomer content (AA), similarly, the contents of the first layer and the second layer were as same as above except keeping that of crosslinking agent (4.5 g), while the third layer of preemulsified monomer mixture of MM A/A A weight ratio was 156.17/4.83, 154.56/6.44, 151.34/9.66, and 149.75/ 11.25, respectively, corresponding to 3, 4, 6, and 7% based on the third layer monomer's weight percent.
Conversion and Particle Size Measurement of the Latexes
At 30-min intervals, samples of the latex (10 mL) were removed into preweighed vials containing 1 mL hydroquinone solution (2 wt%) to prevent the further polymerization; these were surrounded by ice to quench the polymerization and then analyzed gravimetrically to determine the instantaneous conversion (on the basis of the monomer fed until the sampling lime) and overall conversion (on the basis of the monomer fed in the full emulsion polymerization process). Particles sizes were measured with a fixed 90c scattering angle with dynamic light scattering (DLS) on a Malvern Zetasizer (Worcestershire, UK) 3000HS on fine, and the cell temperature was controlled at 25 [+ or -] 0.1[degrees] C. The particle diameters quoted are the mean values of the z-average diameters ([d.sub.z] s) calculated by the cumulate method.
A [PHI] 34 twin screw extruder (L/D = 28, Nanjing Institute of Extrusion Machinery, China) was employed to prepare PA 6/PBMA blends with the weight ratio of 100/20 at a screw speed of 65 rpm and barrel temperatures 220-230-233-223[degrees] C. The pelletized materials were dried and injection molded into standard specimens in an injection-molding machine (JPH-30, Guangdong Hongli Machine, China).
The PBMA latex and PA 6/PBMA blend morphology was examined by TEM (JEM-2I00). The latex was dispersed in water sufficiently with ultrasonic waves before characterization and then prepared by the casting of one drop of diluted solution onto a carbon-coated copper grid. For the blend sample, it was ultra-microtomed in thin films with thickness of 60 nm at -60 [degrees] C to avoid deformation of the dispersed particles.
Impact and Tensile Testing
The mechanical testing was performed at 23 C and 50% relative humidity, and the samples were climatized into this condition for 48 h before testing. The notched Chaipy impact strength was measured with an impact testing machine (ZBC-4) based on National Standard Testing Methods GB 1043-1993, China. The specimen's dimension is 80 mm X 10 mm X 4 mm, notch tip radius 0.25 mm, impact speed 2.9 m/s and energy 4.0 J. The tensile testing was carried out on a universal tensile tester (CMT-6104) according to GB 1040-1992 (equivalent to ASTM D-638). The type of the specimen is V. The specimens dimensions are as follows: length overall (LO) 63.5 mm, gage length (G) 7.62 mm, width of narrow section (W) 3.18 mm, and the thickness (T) 3.2 mm. The tensile strength was measured at a crosshead speed of 50 mm/ min. The average of at least six tests was used.
Dynamic Mechanical Analysis
Dynamic mechanical tests were performed at a frequency of 1Hz by using Tritec-2000 dynamic mechanical analyzer (Triton, UK) under a dual-cantilever beam blending mode. Storage modulus ([pounds sterling]*), loss tangent (tan S) were measured between -80 and 150'C with a heating rate of 3cC/min in a nitrogen atmosphere.
Scanning Electron Microscopy Measurement
The notched Charpy impact-fractured surfaces of PA 6/PBMA blends were observed with a JSM-6490LV scanning electron microscope (SEM). The surface was coated with gold and the accelerating voltage was 10 kV.
RESULTS AND DISCUSSION
Preparation of Poly(BAIPMMA-co-AA) Latexes
Latexes were prepared by seeded emulsion polymerization (i.e., the addition of monomer, initiator, and surfactant to a previously prepared latex), which had the advantage of preventing the uncertainties of the particle initiator stage and, therefore gave better batch-to-batch reproducibility. Instantaneous and overall conversions were calculated from a mass balance of the reagents in the polymerization with the percentage solid content measured at each sampling time:
Instantaneous conversion (%) = (Mass of polymer formed/Mass of monomer added) x 100 (1)
where the mass of monomer added is the sum of the monomer in the seed stage and any monomer that has been added during the growth stage.
Overall conversion (%) = (Mass of polymer formed/Total mass of monomer) x 100 (2)
where the total mass of monomer is the sum of the monomer in the seed stage and all of the monomer in the growth stage.
Plots of conversion versus reaction time for each of the latex preparations were similar. The typical plot of the variation of the conversion versus reaction time with 0.6wt% crosslinking agent (ALMA) and 5 wt% functional monomer (AA) is shown in Fig. la. All of the polymerizations were observed to proceed at high instantaneous conversion (>90%), i.e., most of the monomers added were polymerized. Final conversions were found to be high (>97%) for all of the polymerizations, which showed that a continuation of the polymerization for 1 h after the end of the monomer addition stage was adequate to allow for complete conversion.
The DLS technique was used to obtain quantitative information about the particle sizes of colloidal systems. In this study, DLS provided a rapid means of monitoring the particle size of the latex particles during both the seed and growth stages of polymerization. With this information, it was possible not only to establish and reproduce a latex system of the known particle diameter but also to determine whether, during the growth stage of polymerizations, the latex particles grew sequentially or if the secondary nucleation occurred.
Latex particle diameters were determined by DLS and compared with those theoretically calculated from the following equation:
[d.sup.t] = [([M.sub.t][I.sub.t]/[M.sub.t]).sub.1/3] x [d.sub.s]
where dt is the diameler of the particle at time t, [M.sub.t], is total mass of the monomer added at time t. I, is the instantaneous conversion at time t, [M.sub.t] is mass of monomer added in the seed stage, and [d.sub.s] is the seed particle diameter as measured by DLS.
Plots of particle diameter versus reaction time for each of the latex preparations were similar also. Plot of the particle diameter versus reaction time for the latex with 0.6 wt% ALMA and 5 wt% AA is shown in Fig. lb. Figure 2 shows the distribution of its final particle size poly-dispersity index (PDI). PDI of the latex was 0.035. These results showed that the final particle sizes of the latex presented narrow distribution. The good agreement shown between the experimental and theoretical particle diameters throughout the polymerization for all of the latexes provided strong evidence that the observed particles were grown without significant secondary nucleation and that all of the polymer particles formed were spherical. This, coupled with the low levels of coagulum (<1.70 wt%) measured for all of the latexes, showed that the appropriate contents of the crosslinking agent and functional monomer were used in the growth stage of the polymerization. The high final overall conversions were achieved with a 50% solid content with final particle diameters of 498 nm. The morphology of PBMA latex is observed by TEM, as shown in Fig. 3. It could be seen that the particles consisted of a dark core which was poly (BA) and a brighter shell which was poly (MMA-co-AA) indicated that the core-shell structure of PBMA latex was obvious. From the observation of TEM and DLS measurements, the final PBMA latex was spherical particles with narrow distribution.
Tables 1 and 2 summarized the results of the emulsion polymerization procedure with different crosslinking agent (ALMA) and functional monomer (AA) levels in the second growth stage. The final monomer conversions were all high, and the particle sizes for all of the latexes were almost the same within the experimental error. By the comparison of the final monomer conversions and the measured latex particle sizes, we could see that the latex preparations with different ALMA or AA contents had no significant effect on the final monomer conversion and the particles' formation in the emulsion polymerization procedures of PBMA.
TABLE 1. Summary of the final data for PBMA latexes with different ALMA contents. ALMA The final The final Coagulum Sample content conversion latex particle content No. (wt%) (wt%) size (nm) (wt%) 1 0.5 99.10 505 1.56 2 0.6 97.60 498 1.39 3 0.7 99.20 491 1.62 4 0.8 98.71 494 1.48 5 0.9 97.60 490 1.60 TABLE 2. Summary of the linal data for PBMA latexes with different AA contents. ALMA The final The final Coagulum Sample content conversion latex particle content No. (wt%) (wt%) size (nm) (wt%) 6 3 99.30 494 1.38 7 4 97.73 493 1.68 2 5 97.60 498 1.39 8 6 99.30 494 1.63 9 7 97.65 490 1.41
Blending of PBMA Modifier with PA 6 Resin
TEM Microphotograph of Dispersion of PBMA Particles in PA 6 Matrix. Figure 4 shows TEM image of PA 6/PBMA (Sample 2) blend with the weight ratio of 100/20. The intrinsic microstructure of the blends was prepared from the ultra-microtomed in thin films from the sample. It could be seen that PBMA particles presented a sphericity domain shape in PA 6 matrix, whereas very few were aggregated. A good dispersion of PBMA particles in the matrix was achieved. The particle size of the core-shell modifier, which was set during the synthesis process, could remain after they were dispersed in a host matrix.
Mechanical Properties of PA 6/PBMA Blends. As the WAXD/DSC studies have revealed that the crystallinity of the PA blends and nanocomposites remained unaffected (33), (34), we will just discuss the effect of PBMA particles on the mechanical properties of PA 6 matrix in the following text. It is generally believed that the interfa-cial adhesion between the dispersed rubber particles and the matrix plays an important role in the toughening of polymers. To increase interfacial adhesion and improve the miscibility of polymer blends between matrix and disperse phase, the method of reactive compatibilization is very often used to obtain blends with desirable properties (35), (36). Figures 5 and 6 show the plots of mechanical properties versus crosslinking agent (ALMA) content and functional monomer (AA) content for PA 6/PBMA blends, respectively. According to Wu's classification (37), (38), including entanglement density and characteristic ratio of the chain, PA 6 should be classified as a pseudo ductile polymer. Therefore its notched Charpy impact strength was only 8.50 kJ/[m.sub.2]. The incorporation of PBMA core-shell modifiers greatly improved the notched impact strength and/however an inflexion occurred with the increase of ALMA or AA content.
The notched impact strengths of the PA 6/PBMA blends were mainly influenced by the cross-linking density of the core component, poly (BA), and the interfacial adhesion of the shell component, AA, with the matrix, PA 6. It is believed that the deformation ability of the rubber component was responsible to a great extent for the impact properties of polymer matrix. When ALMA content was below 0.6 wt%, the deformation of rubber particles was not enough to absorb the impact energy, thereby the blends' impact strengths were low (Fig. 5). When ALMA content was over 0.6 wt%, poly (BA) component turned to become rigid particles which lost the toughening ability, resulting in the decrease of the blends' impact strength. The excess tight poly (BA) molecular chains were surely detrimental to the impact toughness of the blends. It was difficult for the hardened rubber partides to produce cavitations' capacity and yielding for the matrix. It was found that using 0.6 wt% ALMA content, the notched impact strength of the PA 6/PBMA blend reached the maximum of 30.3 kJ/[m.sup.2] which increased 256% compared with that of pure PA 6. Only did the cross-linking density for the rubber components reach an optimal value, could the rubber particles play a toughening effect on the PA 6 matrix.
The Charpy impact strength of PA 6/PBMA blends with increasing of functional monomer (AA) content in the shell layer shows the similar trend, as presented in Fig. 6. This was due to fact that at lower mass fraction of AA, the interface between PA 6 and modifiers was weak, which resulted in the modifier particles debonding easily from the PA 6 matrix by external forces, and the propagation of crack through the interface became much easier, which would lead to premature failure in the sample. The interfacial reaction between the carboxyl group of PBMA and the amide functional group of PA 6 matrix during melt processing improved the rubber particles' dispersity and the stress transfer in the blends. The thicker gradient layer around the rubber particles would hinder the separation of rubber particles agglomerate in the matrix polymer. Certain morphological features such as thin filament connecting and even networking of the dispersed rubber phase, as confirmed by SEM observation of the fractured impact surfaces described below may contribute to the overall ductility of the high impact strength of the studied blends. Moreover, the phenomenon of stress concentration caused by the addition of PBMA panicles in the PA 6 matrix could be changed via the rubber particles and PA 6 molecule chains, as well as the effect of interfacial energy transmission. If the interfacial adhesion between the PBMA particles and the PA 6 matrix was suitable, the action of PBMA particles would cause toughening of the PA 6 matrix by initiating shear bands1 Therefore, the ductility of blends became better with increase of the AA content. However, if the interfacial adhesion is too strong, this would cause the decrease the capacity in initiating the shear bands. When 4% AA by weight is used (Fig. 6), the notched impact strength of the blend reached 34.7 kJ/ m2, which increased 308% compared with that of pure PA 6. This led to the significant improvement of the impact toughness of the blends, and was considered to explain the toughening mechanism.
The tensile strength is an important characteristic of polymeric materials because it indicates the limit of final stress in applications. As compared with pure PA 6, with increase of the ALMA or A A contents, the blends' tensile properties showed a slow decrease (Figs. 5 and 6). This is due to the fact that the rubber component had low modulus and tensile strength. The smaller decrease represents the existence of an interface region by which two phases are bonded strongly because the tensile strength is strongly dependent on the interface structure of the blends. The stress has to transfer across the interface to avoid the fracture. Noolandi, Hong previously showed that the thickness of the interface between A and B polymer phases increased with the addition of A-B block or graft polymers (39). Strong interactions result in good adhesion and efficient stress transfer from the continuous to the dispersed polymer phase in the blends. Therefore, an improved toughness-stiffness balance may be obtained in PA 6/PBMA blends by introduction of reactive compatibility between the core-shell polymer and the matrix, PA 6, and the suitable cross-linking density for the core component, poly (BA).
SEM Microphotographs of Fractured Surface of the Blends. Figure 7 shows SEM micrographs of the impact fractured surfaces of PA 6/PBMA blends with different AA contents. The surface of pure PA 6 (Fig. 7a) was rough with many topographical irregularities indicating the homogeneous deformation of the matrix phase. With the increase in AA contents, the voiding is obvious because the functional monomer (AA) was reacted with the end amide group of PA 6 (Fig. 7b-d). When there are sufficient reactive groups on the surface of PA 6/PBMA blends, a good compatibilization and dispersion can be achieved. It is expected that the compatibilization reactions at the interface will lead to a decrease in the interfacial tension and some stabilization against coagulation. However, in Fig, 7c, the morphology of the optimized blend matrix (5 wt% AA content) reveals evenly distributed fibrillation network structures originating from within the bulk of the material. These fiber Haled network structures indicate the discontinuous PBMA-phase which was added to the PA-6 matrix for toughness enhancement and offered resistance to the fracture process, via energy dissipat ion mechanism, since poly (BA) is a softer matter with much lower [T.sub.g], as compared with that o(PA-6. Thus it may be well comprehended that with the increase in the AA content the mechanism of PA 6-phase deformation undergoes a ductile transition behavior. These photomicrographs are quite consistent with the increase in toughness observed by adding PBMA modifier to the PA 6 matrix (Fig. 7c). On the basis of the above results, it could be concluded that the interfacial adhesion and morphological features are important for toughening PA 6 and this material shows wide practical uses in industry.
Dynamic Mechanical Properties of PA 6/PBMA Blends. Dynamic mechanical analysis (DMA) is a powerful method to study polymer miscibility. To examine the miscibility of PA 6/PBMA blends, a series of PBMA modifiers were prepared by semicontinuous emulsion polymerization with different cross-linking agent (ALMA) contents in the core layer and different functional monomer (AA) contents in the shell layers. The mechanical relaxation process of the pure PA 6 and PA 6/ PBMA blends were measured between -100 and 150[degrees]C. The curves in Fig. 8a and b, obtained by DMA measurement, show the variation of storage modulus and loss tangent for PA 6 and for the PA 6/PBMA blends as a function of temperature with different ALMA contents of PBMA component. It is well known that PA 6 has two relaxation peaks: an a peak at 51.4[degrees] C and a [beta] peak at -71.0[degrees]C. The a peak is believed to be associated with the glass transition ([T.sub.g] PA 6), the [beta] peak is due to segmental motion of the amide groups which are not bonded to other amide groups ([T.sub.[beta]]) (23). As seen from Fig. 8, the glass transition peak of poly (BA) (Tg PBA) in PBMA component also appears at - 44.4[degrees] C. Meanwhile, the low temperature transition for PA 6 and PBMA showed a tendency to merge together from -71.0[degrees] C to - 44.4[degrees] C, with one peak forming a shoulder on the other. A broadened transition was noted, which meant the interaction between the PA matrix and PBMA dispersed phase. The breadths of low temperature transition zones ([T.sub.g] PBA-[T.sub.[beta]] PA 6) were 17.5, 26.2, 26.2, 22.8, 17.5[degrees]C for PA 6/PBMA blends with ALMA contents of 0.5, 0.6, 0,7, 0.8, 0.9 wt% in PBMA, respectively. These results indicated that the transition zone became broadened with the increase in cross-linking agent content. The same results could be found at the glass transition region of PA 6 after the addition of PBMA component, formed a continuous transition region from 51.4[degrees]C to 86.5[degrees]C, which is the glass transition temperature of the PMMA shell phase. This evidence for large-scale phase separation or heterogeneity was not apparent. The enhancement in the tan S due to ALMA has been observed to be clearly distinguishable, i.e., the sample 2 with 0.6 wt% of ALMA content has shown a dual-peak distribution over a broader temperature regime than the other blends, such as the ones with 0.5, 0.7, and 0.8 wt% of ALMA, i.e., samples 1, 3, and 4. These results was caused by the improvement at the interface between the PBMA component and PA 6 matrix and showed that PA 6 was partially miscible with the PBMA phase. Figure 9 presents the DMA curves of PA 6/PBMA blends with different AA contents in PBMA component. The low temperature transition became broader, and the values of [T.ub.g] PBA-T[beta] PA 6 were almost the same for different AA contents. The high glass transition regions were broadened for all samples. The striking differences came from the magnitude of the loss peak. The value of tan & increased with increasing the AA content in PBMA phase. For the sample 7 (4 wt% AA content), the value of tan & is 0.17 because of the interfacial reaction between the carboxyl group of AA component and the amide group of PA 6 during melt extrusion. When the AA content was over 4 wt% (Sample 9), the magnitude of tan S at high temperature transition was very low, - 0.14. This result was consistent with the mechanical properties. It was well documented (40), (41) that the resistance to impact required a dynamic mechanical energy dissipation mechanism at the temperature and frequency of the impact. It is apparent that the addition of the PBMA leads to the appearance of two merged loss peaks in the range--71.0 to -44.4 C and 51.4 to 86.5[degree] at the DMA curves of PA 6/PBMA blends and the value of the high transition loss peak increases with ALMA or AA content in PBMA component, implying that the impact strength of the PA 6 blends should increase with cross-linking density for the core layer and functional group, AA for the shell layer of core-shell structured polymer, PBMA. Thus, the improvement of the impact toughness depends not only on the breadth of the transition zone, but also on the magnitude of loss peak.
Acrylic core-shell panicles, poly(n-butyl acrylate)/poly(methyl methacrylate-co-acrylic acid), i.e., poly(BA/MMA-co-AA) (PBMA), with narrow size distributions were successfully prepared in monomer-starved semicontinuous emulsion polymerization. The synthesized PBMA core-shell latexes were used to toughen PA 6 matrix. Selection of reaction conditions makes possible the preparation of structured core-shell polymers that contain rubber component as core and functional component as shell on the surface and obtains sufficient dispersion. The PA 6/PBMA blends had significant toughening effect because of the good dispersion and good adhesive interface between PA 6 and PBMA. The cross-linking agent (ALMA) content in the core layer and the functional monomer (AA) content in the core layer also influenced the mechanical properties of PA 6/PBMA blends. The fractured surfaces of the blends exhibited significant ductile deformation of the PA 6 matrix, which was consistent with the impact properties of the blends. Depending on the type of matrix, the amount of carboxyl groups on the surface of the core-shell particles can be regulated during the polymerization. This method of preparing functionalized core-shell particles is very promising in achieving well compatibilized thermoplastics.
The authors thank Prof. Peter A. Lovell (University of Manchester, UK) for many fruitful discussions.
(1.) M.J. Modic and L.A. Pottick. Polym. Eng. Set., 33. 819 (1993).
(2.) C.J. Wu, J.K. Kuo, and C.Y. Chen, Polym, Eng. Set, H. 1329 (1993).
(3.) M. Lu. 11. Kcskkula. and D.R. Paul. Polym, Eng. Sci.. 34. 33(1994).
(4.) G. Burgish, M. Paternoster. N. Pcduto, and A. Saraceno, J. Appl. Polym. Sci.. 66, 777 (1997).
(5.) T. Harada, E. Carone Jr.. R.A. Kudva, H. Kcskkula, and D.R. Paul, Polymer, 40, 3957 (1999).
(6.) D.M. Laura, H. Keskkula, J.W. Barlow, and D.R. Paul, Polymer. 44, 3347 (2003).
(7.) J.J. Huang, H. Keskkula, and D.R. Paul, Polymer, 45, 4203 (2004).
(8.) O. Okada, H. Keskkula, and D.R. Paul, J. Appl. Polym. Sci.. 42, 1739 (2004).
(9.) I. Kelnar, J. Kotekal, and B.S. Kapralkova, J. Appl. Polym. Sci.. 96, 288 (2005).
(10.) J.J. Huang, H. Keskkula, and D.R. Paul, Polymer, 47, 639 (2006).
(11.) S.C. George, G. Groeninckx, K.N. Ninan, and S. Thomas, J. Polym. Sci B, 38, 2136 (2000).
(12.) P. Van Puyvclde, Z. Oommen. P. Kocts. G. Groeninckx. and P. Moldcnaers, Polym. Eng. Sci.. 43, 71 (2003).
(13.) E. Radovanovic, E. Carone Jr., and M.C. Goncalves, Polym. Test., 23, 23 I (2004).
(14.) S.H. Jafari. P. Potschke, M. Stcphan. and H. Alberts. Polymer, 43, 6985 (2002).
(15.) S.L. Sun, Z.L. Tan, X.F. Xu, C. Zhou, Y.H. Ao, and H.X. Zhang, J. Appl. Polym. ScL, 43, 2170 (2005).
(16.) K. Premohet-Sirisinha and S. Chalcarmthitipa, Polym. Eng. 5c/.. 43.317 (2003).
(17.) J.J. Huang and D.R. Paul, Polymer, 47, 3505 (2006).
(18.) XII. Wang. H.X. Zhang, W. Jiang, Z.G. Wang, C.H. Liu, H.J. Liang, and B.Z. Jiang. Polymer, 39, 2697 (1998).
(19.) I. Vieira, V.L.S. Severgnini, D.J. Mazcra, M.S. Soldi, and E.A. Pinhciro, Polym. Degrad. Stab., 74, 151 (2001).
(20.) C. Wang. J.X. Su. J. Li. H. Yang, Q. Zhang. R.N. Du. and Q. Fu. Polymer. 47, 3197 (2006).
(21.) M. Lu, H. Kcskkula, and D.R. Paul, J. Appl. Polym. Sci., 58, 1175 (1995).
(22.) M. Lu. H. Kcskkula, and D.R. Paul, J. Appl. Polym. Sci., 59, 1467 (1996).
(23.) Z.Z. Yu, Y.C. Ou, Z.N. Qi, and G.H. I lu, J. Appl. Polym. Sci., 36, 1987 (1998).
(24.) X. Wang, H. Li, and E. Ruckcnslein. Polymer. 42. 9211 (2001).
(25.) A. Valcnza, A.M. Visco. and D. Acicrno, Polym. Test., 21, 101 (2002).
(26.) H. Janik, R.J. Gaymans. and K. Dijkstra. Polymer. 36, 4203 (1995).
(27.) E. Caronc Jr., U. Kopacak, M.C. Goncalves, and S.P. Nuncs. Polymer. 41, 5929 (2000).
(28.) J. Pong, J. qiao, S. Zhang, and G. Wei. Macromol. Maier. Eng., 287, 867 (2002).
(29.) C.R. Kumar, S.V. Nair, K.E. George, Z.O. Ommcn, and S. Thomas, Polym. Eng. Sci.. 43. 1555 (2003).
(30.) R. Chowdhury. M.S. Bancrji, and K. Shivakumar. J. Appl. Polym. Sci., 104, 372 (2007).
(31.) X. Ding, R. Xu, D. Yu, H. Chen, and R. Fun. J. Appl. Polym. Sci., 90. 3503 (2003).
(32.) X. Zhang, Y. Liu, J. Gao, F. Huang, Z. Song, G. Wei, and J. Qiao, Polymer. 45, 6959 (2004).
(33.) C. Wang, J.X. Su. J. Li. H. Yang. Q. Zhang. R.N. Du. and Q. Fu, Polymer. 47, 3197 (2006).
(34.) N. Dayma and B.K. Satapalhy, Mater. Des., 31, 4693 (2010).
(35.) Y.J. Sun, G.H. Hu, M. Lambla, and H.K. Kotlar, Polymer, 37.4119 (1996).
(36.) G.H. Hu, Y.J. Sun, and M. Lambla, Polym. Eng. Sci., 36, 676 (1996).
(37.) S. Wu, Polym. Eng. Sci., 30, 753 (1990).
(38.) S. Wu, Polym. Int., 29. 229 (1992).
(39.) J. Noolandi and K.M. Hong, Macromolecules, 17, 1531 (1984).
(40.) E. Sachcr. Polymer. 21. 1234 (1980).
(41.) E. Sacher, J. Macromol. Sci., B15. 257 (1978).
Haiyan Zhao, (1) Yanmei Yao, (1) Xinran Liu, (1) Qingxin Zhang, (1) Nongyue Wang, (1) Liqun Zhang, (2) Xiongwei Qu (1)
(1.) Institute of Polymer Science and Engineering, School of Chemical Engineering, Hebei University of Technology, Tianjin 300130, People's Republic of China
(2.) Beijing University of Chemical Technology, Beijing 100029, People's Republic of China
Correspondence to: Xiongwei Qu; e-mail: firstname.lastname@example.org
Contract grant sponsor: Natural Science Foundation of Hebci Province; contract grant number: E2010000107; contract grant sponsor: Key Lab of Beijing City on Preparation and Processing of Novel Polymer Materials; contract grant number: 2006-1: contract grant sponsor: Distinguished Youth Scientist of NSF; contract grant number: 50725310.
Published online in Wiley Online Library (wileyonlinelibrary.com).
[c] 2011 Society of Plaslies Engineers
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|Author:||Zhao, Haiyan; Yao, Yanmei; Liu, Xinran; Zhang, Qingxin; Wang, Nongyue; Zhang, Liqun; Qu, Xiongwei|
|Publication:||Polymer Engineering and Science|
|Date:||Feb 1, 2012|
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