Article pubs.acs.org/Macromolecules
Phase Separation of Poly(methyl methacrylate)/Poly(styrene-coacrylonitrile) Blends with Controlled Distribution of Silica Nanoparticles Chongwen Huang, Jianping Gao, Wei Yu,* and Chixing Zhou Advanced Rheology Institute, Department of Polymer Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, P. R. China S Supporting Information *
ABSTRACT: Effects of selective location of silica nanoparticles on the phase separation of poly(methyl methacrylate)/poly(styrene-coacrylonitrile) (PMMA/SAN) blends were investigated via combination of rheological method and optical microscopy. Through grafting polystyrene chain to the surface of silica nanoparticles, the silica nanoparticles were controlled to selectively locate at interfaces or in the PMMA-rich domains. Power-law analysis of the moduli and shifted Cole−Cole plots were applied to determine rheological transition temperature (apparent binodal temperature) of blend with near-critical and off-critical compositions for both neat blends and particle-filled blends. The particle location had significant influence on the rheological transition temperature but little impact on optically determined binodal temperature. This discrepancy was discussed through morphology observation via transmission electron microscopy (TEM) for blends under different phase separation conditions. It was found that nanoparticles retard coarsening of morphology during phase separation. The most striking slowdown was found in off-critical blends with nanoparticles located on the interface. On the other hand, nanoparticles preferentially locating in the minor phase could act as nucleation sites but decreased the total number of nuclei. The difference in the rheological transition temperatures is ascribed to the effect of nanoparticles on the components’ viscoelasticity and the morphology during phase separation. superposition (TTS) principle9,14,20,21 and Cole−Cole plot (plotting η″ vs η′)14,20 are most frequently used to distinguish whether the blends are in the homogeneous or heterogeneous state. These methods utilize the dynamic viscoelastic data at various temperatures, and the accuracy of the transition temperature thus determined is affected by the width of the temperature range examined. On the other hand, temperature ramp has been also widely used to determine the phase transition temperature.5,14,15,22−24 With increasing temperature, the storage modulus G′ usually decreases owing to the enhanced chain mobility. However, when the blend enters transition region, the decreases in G′ can be compensated by an increase due to contribution from interface and concentration fluctuation, thereby slowing this decrease and sometimes even increasing the storage modulus. For example, in polybutadiene (PB)/low vinyl content polyisoprene (LPI) blends5 or in polystyrene (PS)/poly(vinyl methyl ether) (PVME) blends4 having a large dynamic asymmetry of the components, an evident increase in G′ has been observed. For polymer blends with weaker dynamic asymmetry,19 the enhancement of G′ due
1. INTRODUCTION Polymer blends have been the subject of research for decades owing to their wide applications. As a key factor, the phase diagram or phase separation temperature of polymer blends has attracted considerable attention.1−5 Optical microscopy6,7 and light scattering6,8−10 have been frequently used to obtain the phase diagram or phase separation temperature. Obviously, the transparency and thickness of sample would inevitably affect the application of optical or light scattering method, which makes these methods unsuitable for blends of components having similar refractive indices or blends filled with high load of solid particles.11 At the same time, rheological measurements have been widely used to detect the phase separation4,5,12−16 because rheology can link the viscoelastic response of polymer blends to subtle structural changes during phase transition regardless of the transparency and refractive indices of the components. Specifically, the viscoelasticity of polymer blends in the terminal region can be directly associated with the interfacial tension and characteristic length of domains.17,18 It is usually believed that higher resolution can be achieved through the rheological measurements at various oscillation frequencies than the optical measurements.19 Various protocols have been proposed to rheologically detect the phase separation temperature. The time−temperature © XXXX American Chemical Society
Received: June 12, 2012 Revised: September 20, 2012
A
dx.doi.org/10.1021/ma301186b | Macromolecules XXXX, XXX, XXX−XXX
Macromolecules
Article
to the interface and concentration fluctuation is relatively small but still detectable. All these experiments show that the rheological detection of phase separation is rather complicated, which is ascribed to relative contribution of concentration fluctuation to the moduli25 as well as the competition between the time scales of rheological measurement and phase separation. Actually, direct comparisons between the rheological methods and other methods like optical method should be cautious since their spatial resolutions are generally different. However, a detailed evaluation on such difference by considering the viscoelastic mismatch in components, the kinetics of phase separation, and the relative strength of concentration fluctuation is still missing even for binary blends, not to mention more complicated particle-filled polymer blends. Over the past decades, incorporation of nanoparticles into polymer blends has also attracted considerable attention because the nanoparticles can be used to control the morphology of polymer blends,26−33 act as the compatibilizers for immiscible polymer blends,34,35 and change the phase separation temperature.36−42 Recently, the influence of nanoparticles on the phase behavior of polymer blends has gradually become the concern of investigations. Slowdown of the coarsening during phase separation due to the nanoparticles is often observed,41 especially when the particles are located at the interface between different phases. However, the effect of nanoparticles on the phase separation temperature is not well understood. Lipatov et al.37,38 found that the silica particles significantly shifted the phase diagram of chlorinated polyethylene (CPE)/copolymer of ethylene with vinyl acetate (EVA) and PVA/PMMA blends due to considerable compatibility effect, but the particle location was not specified. Recently, the effect of silica particles selectively located in PMMA phase on the phase behavior of PMMA/SAN was investigated via the rheological method.39 An increase in rheological transition temperature (apparent binodal temperature) was observed for all compositions, and the shift of phase diagram was much larger when PMMA formed the minor phase. Gharachorlou et al.43 showed the increase of binodal temperature in PS/PVME (polystyrene/poly(vinyl methyl ether)) blends on addition of silica nanoparticles using rheology, turbidity, and DSC methods. No data were shown for the effect of nanoparticles on the rheological properties of components although silica particles were believed to locate in PVME phase. The effect of layered silicates on the phase behavior and morphology was investigated by Yurekli et al.,40,41 who found that the phase diagram was almost unaffected by the silicate with the volume fraction up to 0.008 and that the morphology depended on the size of layered silicates. In a theoretical view, Ginzburg42 proposed a simple model to investigate the influence on thermodynamics of polymer blend from nanoparticles being selectively dispersed in one of the polymeric components. However, this theory needs more comparisons with experiments. In our previous work, we found that the selective segregation of silica particles into PMMA enhances the phase stability when PMMA forms the minor phase, whereas this influence on the phase behavior is nearly negligible when PMMA forms the major phase.44 Despite those intensive studies, an overview of the effect of the particles location on the phase diagram and phase separation kinetics still remains unclear. Moreover, as a promising method to study the phase behavior of polymer blends with high load of solid particle, the rheological method needs to be further examined.
In this work, we focus on the blends of the poly(methyl methacrylate)/poly(styrene-co-acrylonitrile) (PMMA/SAN). The location of nanosilica (SiO2) particles is controlled by varying the chain length of grafted polymers through surfaceinitiated ATRP. Power-law analysis of the moduli and Cole− Cole plots are applied to obtain rheological transition temperature (apparent binodal temperature) of blends with and without nanoparticles. Optical microscopy is used to detect the thermodynamic phase diagram of neat polymer blends and blends filled with different nanoparticles. The discrepancy of phase diagrams obtained with these two methods is discussed. The influence of the selective location of silica particles on the phase behavior is also investigated by combination of rheology and morphology observation.
2. EXPERIMENTAL SECTION 2.1. Materials. Poly(methyl methacrylate) (PMMA, IF850, Mw = 15.9 × 104 g/mol, Mw/Mn = 1.64, Tg = 96 °C) and poly(styrene-coacrylonitrile) (SAN, 81HF, Mw = 14.1 × 104 g/mol, Mw/Mn = 2.08, Tg = 105 °C, containing 28.4 wt % acrylonitrile) were both supplied by LG Chemical Ltd. Hydrophilic fumed silica nanoparticles with specific surface area 160 ± 20 m2/g and average particle diameter 30 nm (data provided by the supplier), denoted as SiO2−OH, was purchased from Hang Zhou Wan Jing New Material Co. Ltd. Details of synthesis and characterization of the PS chains are shown in the Supporting Information. Silica nanoparticles grafted with short, medium, and long polystyrene brushes are denoted as SiO2−PS, SiO2−PSm, and SiO2− PSl, respectively. 2.2. Sample Preparation. PMMA, SAN, and various silica nanoparticles were all dried at 80 °C in a vacuum oven for at least 48 h before processing. The PMMA/SAN/nanoparticle blends were denoted by A/B/x, where A and B represent the weight fraction of PMMA and SAN in the binary blend, respectively, and x is the weight fraction of nanoparticles with respect to the total amount of polymers. After melt blending in a torque rheometer (XSS-300, Shanghai Kechuang Rubber & Plastic Equipment Co, China) at 150 °C and 50 rpm for 20 min, disk-shaped specimen (of 25 mm diameter and 1 mm thickness) for rheological tests were prepared by compression molding at 150 °C under 10 MPa. Since the (LCST-type) phase separation temperature was well above 150 °C, the as-mixed samples were transparent and homogeneous. 2.3. Instrumental Analysis. The viscoelastic properties of the blends were measured with a stress controlled rotational rheometer (Bolin Gemini 200HR, Malvern Instrument, UK). The parallel-plate geometry with plate diameter 25 mm was used in all the experiments, and the gap between the plates was 1 mm. Isothermal dynamic frequency sweep experiments were carried out at various temperatures ranging from 150 to 200 °C. The strain in all measurements was kept as 2% to ensure the linearity of response. To characterize the morphology of these nanocomposites by transmission electron microscopy (TEM), samples were microtomed with a diamond knife at room temperature. The TEM images were obtained by a JEOL JEM-2100 instrument operating at an accelerating voltage of 200 kV. To enhance the contrast between the components in the blends, SAN phase was selectively stained by RuO4 for about 10 min. Optical observation was conducted on a Leica optical microscopy (DM2500P), and images were collected by a CCD camera with 15 frames/s. The experimental temperature was controlled by a Cambridge Linkam shear cell (CSS450). Samples were compressed at 150 °C into about 200 μm in thickness before observation. The temperature ramp was first applied to monitor the morphology evolution of polymer blends at a heating rate of 1 K/min from 160 to 200 °C, and the gray scale of recorded image throughout the ramp was calculated and converted to transmission. In order to obtain the static phase diagram of polymer blends, the samples were annealed at several different temperatures for 2 h. The gray scale of recorded image was also monitored during annealing and converted to transmission. B
dx.doi.org/10.1021/ma301186b | Macromolecules XXXX, XXX, XXX−XXX
Macromolecules
Article
3. RESULTS AND DISCUSSION 3.1. Selective Distribution of Silica Particles. Recently, Chung et al.45 have observed the selective segregation of PMMA-grafted nanosilica fillers at the interfaces in PMMA/ SAN blends. In the present contribution, the modification of surfaces of silica particles has been conducted via surfaceinitiated atom transfer radical polymerization of styrene. After grafting PS to the particle surface of silica, the dispersion of nanoparticles in the PMMA/SAN blends was analyzed. Figure 1 shows the TEM image of SiO2−PS filled blend annealed at
explained in the Supporting Information) are preferentially located at the interface, as noted in Figure 2b. (This was the case also for the SiO2−PS particles having shorter PS grafts of M = 500 g/mol; cf. Figure 1.) Nevertheless, the selective localization in the PMMA-rich domains occurs again for the SiO2−PSl particles having longer PS grafts (M = 46 000 g/ mol), as seen in Figure 2c. In respective cases, the particles tend to aggregate with each other. One possible reason for this localization of the SiO2−PSl particles can be found from the interfacial tension α between components: αPS/PMMA = 0.8− 1.02 mN/m for PS and PMMA at ∼200 °C,46,47 whereas αPS/SAN for PS and SAN at 200 °C increases from 0.35 to 1.3 mN/m as the AN content in SAN is increased from 5.4 to 16.8 wt %.48 From the increase of αPS/SAN with the AN content, αPS/SAN in our system is estimated to be 1.8 mN/m, which is larger than αPS/PMMA. Thus, for SiO2−PSl particles possibly having full surface coverage with the long PS grafts and behaving as PS on contact with PMMA and SAN,45 localization would occur in the PMMA domains rather than the SAN domains, as observed in Figure 2c. 3.2. Apparent Phase Diagram from Rheology. Rheology serves as an efficient tool to detect the subtle viscoelastic change in polymer blends and has been widely used to infer the (apparent) phase separation temperature.4,5,12−16 The time−temperature superposition (TTS) principle was frequently used to determine this temperature, on the basis of an assumption that the blends in homogeneous state are thermorheologically simple. Before testing the TTS principle to blends containing particles, it is necessary to test whether it is valid for respective components filled with the particles. As shown in Figures 3a and 3b, TTS works quite well for the component PMMA and SAN containing the SiO2−PS particles (having PS grafts of M = 500 g/mol), especially at low frequencies. Delicate deviations from TTS seen at higher frequencies do not affect the judgment of the phase separation using low-frequency data. For blends containing the particles, TTS works well at low temperatures T to confirm the one-phase structure of the components. An example is shown in Figure 3c for a typical particle-filled system, PMMA/SAN/SiO2−PS obeying TTS at T ≤ 170 °C. In contrast, failure of TTS due to phase separation is clearly observed at higher T, for example, at T ≥ 180 °C in Figure 3c. In general, there are two different mechanisms of phase separation depending on the compositions. Blends with nearcritical compositions experience spinodal decomposition to form cocontinuous morphology in the initial stage of phase separation. In contrast, blends with off-critical compositions undergo nucleation and growth process, first forming nuclei
Figure 1. TEM image of PMMA/SAN/SiO2−PS 50/50/3 annealed at 200 °C for 12 h.
200 °C for 12 h. It is clear that the SiO2−PS particles mainly segregate to the interfaces between unstained PMMA and stained SAN domains (white and dark regions). It has been reported that the location of silica particles (either at the interface or in the PMMA phase) can be controlled via tuning the end group (Cl) or the length of the PMMA grafts on the particles.45 In contrast, our nanoparticles have PS grafts incompatible with either SAN or PMMA. Thus, the thermodynamically stable state could be achieved when those nanoparticles preferentially segregate to the interfaces. Bare SiO2−OH nanoparticles were selectively located in PMMA-rich domains after 2 h annealing, as seen in Figure 2a (and also noted in previous AFM observation44). In contrast, the SiO2−PSm particles having the PS grafts of the chain length M = 2000 g/mol (with M being determined with a method
Figure 2. TEM images of (a) PMMA/SAN/SiO2−OH 70/30/3 annealed at 172.5 °C for 2 h, (b) PMMA/SAN/SiO2−PSm 50/50/3 annealed at 200 °C for 12 h, and (c) PMMA/SAN/SiO2−PSl 80/20/3 annealed at 200 °C for 12 h. PMMA-rich and SAN-rich domains are observed as white and dark regions, respectively. C
dx.doi.org/10.1021/ma301186b | Macromolecules XXXX, XXX, XXX−XXX
Macromolecules
Article
different frequencies converge in a narrow range of T, and the arithmetic mean in this range, 177 and 175 °C for the PMMA/ SAN/SiO2−PS systems with 80/20/3 and 70/30/3 compositions, can be utilized as the rheological transition temperature. It should be emphasized that a percolated particle network structure was formed in none of the PMMA/SAN/SiO2−PS systems examined because of the low particle loading. This fact guarantees that the critical gel-like behavior (frequencyinsensitive tan δ) is exclusively related to the cocontinuous network-like morphology formed by the component polymers. For phase-separated off-critical blends, the storage modulus curve at low frequencies exhibited a shoulder attributable to relaxation of the droplets formed therein, and no critical gel-like behavior (frequency-insensitive tan δ that corresponds to power-law scaling of the moduli) was observed. Thus, Cole− Cole plot was applied to estimate the rheological transition temperature. As an example, Figure 5 shows the shifted Cole− Figure 3. Time−temperature superposition of polymer/nanoparticles composites (a) PMMA/SiO2−PS 100/3, (b) SAN/SiO2−PS 100/3, and (c) PMMA/SAN/SiO2−PS 70/30/3. The reference temperature (Tref) is 160 °C for all the nanocomposites.
and then growing up into droplet-matrix morphology. The cocontinuous morphology formed in the near-critical blends resembles the gel network structure, so that the criteria used for critical gelation can also be used to determine the rheological transition temperature (apparent binodal temperature) of the blends. Specifically, tan δ for critical gel is independent of frequency when the frequency is lower than a critical value.49 Thus, in the plots of tan δ of the blends at low frequencies against temperature, the intersection of the plots for different frequencies can be used to define the rheological transition temperature. In reality, the plots include experimental errors, so that the arithmetic mean of the intersections is used as the rheological transition temperature.44 This method was originally suggested for pure polymer blends but is applicable also to the PMMA/SAN/SiO2−OH system.44 The same method was found to be applicable to near-critical PMMA/SAN/SiO2−PS systems examined in this study, as demonstrated in Figure 4. At lower frequencies, the tan δ at
Figure 5. Shifted Cole−Cole plot for blends for off-critical PMMA/ SAN/SiO2−PS 40/60/3 blend.
Cole plots of η″/aT against η′/aT at different temperatures for PMMA/SAN/SiO2−PS 40/60/3, where aT is the horizontal shift factor in TTS. The plot exhibits one circular arc at low T, but a tail of the plots emerges at high T > 177 °C. This tail grows into another circular arc (representing the interfacial relaxation of the droplet) at sufficiently high T. Thus, the rheological transition temperature can be estimated as the temperature where the tail in the Cole−Cole plots emerges, ∼176 °C for the particle-filled blend examined in Figure 5. With the two methods explained above, the apparent phase diagram of PMMA/SAN/SiO2−PS was obtained. The results are shown in Figure 6. For comparison, the phase diagram of PMMA/SAN neat blend and those filled with SiO2−OH are also shown.44 The SiO2−OH particles are selectively dispersed into the PMMA-rich domains, while SiO2−PS particles were preferentially located at the PMMA/SAN interfaces. Comparison of rheologically determined phase diagrams for the neat and SiO2−OH particle-filled blends (unfilled square and circle) suggests that the particles enhance the apparent stability of the blend to increase the rheological transition temperature (apparent binodal temperature) when the PMMA-rich phase containing those particles is the minor phase, but this enhancement vanishes when PMMA forms the major phase. In contrast, SiO2−PS particles localized at the PMMA/SAN interface always enhance this stability. This enhancement, with
Figure 4. Loss tangent obtained from dynamic frequency sweep at different temperatures for near-critical compositions of PMMA/SAN/ SiO2−PS blend: (a) 80/20/3; (b) 70/30/3. D
dx.doi.org/10.1021/ma301186b | Macromolecules XXXX, XXX, XXX−XXX
Macromolecules
Article
Figure 6. Phase diagram of PMMA/SAN, PMMA/SAN/SiO2−OH, and PMMA/SAN/SiO2−PS blends obtained via rheology and optical microscopy. The content of nanoparticles is 3 wt % in the latter two blends.
respect to the neat blends, is more significant when PMMA is the major component therein. It should be emphasized that the above stabilization effect is different for the particles localized at the interface and in the PMMA-rich phase. In near-critical blends such as PMMA/SAN 70/30, the former type of the particles considerably increases the rheological transition temperature whereas the latter type of the particles hardly changes this temperature. The opposite trend is noted in off-critical blends such as PMMA/SAN 70/30. This difference can be related to the interaction between the particles and component polymers that would change according to the particle location. This point is further discussed later in section 3.4. 3.3. Phase Diagram from Optical Microscopy. Optical methods, including small-angle light scattering and optical microscopy (OM), have been often used to detect the phase structure. Obviously, the transmission of light depends on the transparency of the samples. Thus, optical method is sometimes unsuitable, especially in blends filled with concentrated nanoparticles. Fortunately, in our PMMA/SAN blends, both components have excellent transparency and the fraction of silica nanoparticles is rather small, so that the change of transmittance can be easily monitored to characterize the phase separation process. Figure 7a shows typical OM images of PMMA/SAN 70/30 blend at different temperatures during heating at a rate of 1 K/ min. The sample becomes darker; i.e., its transmittance decreases as the temperature increases. However, no structure can be observed from the OM image (even after its Fourier transformation). The refractive indices of PMMA and SAN are about 1.4914 and 1.567, respectively.50 Although the difference of these indices is not very large, we can still observe the phase separated morphology by OM at higher temperature in the late stage of phase separation. Thus, the change of the light transmittance (change of the grayscale in the OM image) seen in Figure 7a can be ascribed to the scattering from phaseseparated domains of the sizes smaller than the wavelength of light λ. (Note that the domain size in PMMA/SAN system during phase separation is about tens of nanometers, and the domains coarsen to micrometer-sized only after long time annealing, as shown in Figure 2.)
Figure 7. (a) Optical microscopic images of PMMA/SAN 70/30 blend at different temperatures during heating at a rate 1 K/min. (b) Temperature dependence of transmittance for PMMA/SAN neat blends.
The relative variation in the grayscale of the OM image gives a change of the light transmittance. For consistency, the initial gray scales of all samples (corresponding to no phase separation) were normalized to unity. For PMMA/SAN neat blends, the temperature dependence of the transmittance after this normalization is shown in Figure 7b. The transmittance decreases with increasing temperature T > 170 °C, which reflects the phase separation (at a length scale < λ). Nearcritical blends (with the composition of 80/20 and 70/30) exhibit similar magnitudes of the transmittance reduction, while off-critical blends (50/50 and 40/60) show significantly weaker reduction. This difference between the near-critical and offcritical blends is ascribed to a difference of the phase separation mechanisms therein. Near-critical blends experience spinodal decomposition to form cocontinuous morphology that could scatter the light more significantly compared to the droplet morphology formed through nucleation and growth in offcritical blends. For neat blends, we adopted the temperature at which the transmittance shown in Figure 7b decreases by 1% (99% transmittance) as the nonisothermal phase separation temperature (apparent cloud point under nonisothermal condition). This apparent cloud point of the neat blends, shown in Figure 6 with star symbol, is close to the rheological transition points (unfilled square) when PMMA is the major component but becomes higher than the rheological transition point as the SAN content increases. As will be shown below, none of the rheological transition point and optically detected apparent cloud point is identical to the phase transition temperature at equilibrium, and kinetic factors play quite important roles. In order to eliminate the kinetic factors in nonisothermal experiments like temperature ramp, we attempted to examine the phase separation under isothermal condition. Figure 8 E
dx.doi.org/10.1021/ma301186b | Macromolecules XXXX, XXX, XXX−XXX
Macromolecules
Article
isothermal annealing at 154 °C for 2 h but exhibits rapid decay at higher temperatures. Note also that the phase separation is still slow even at the highest T (= 162 °C) examined in Figure 8a so that the micrometer-sized domains was not observed with optical microscopy, as explained earlier. Off-critical PMMA/SAN 40/60 (right panel of Figure 8a) exhibits qualitatively similar behavior, but the decay of transmittance is less significant because of its droplet-matrix morphology (that exhibits weaker scattering compared to the cocontinuous morphology in the near-critical blends). For neat blends, we adopted the same criterion (1% reduction of transmittance) for the isothermal cloud point as utilized in the temperature ramp test, although this criterion might somehow overestimate the isothermal cloud point for off-critical blends of LCST type (including our PMMA/SAN). With this criterion, the cloud points for PMMA/SAN 70/30 and 40/60 blends shown in Figure 8a are estimated to be 155 and 164 °C, respectively. For nanoparticles filled polymer blends, one must be cautious when using optical method to judge the phase separation because the nanoparticles affect the transparencies of samples. It is expected that localization of nanoparticles at the interface or in one phase could enhance the light scattering and reduce the light transmittance, which means smaller domains could be detected with the help of nanoparticles. The evolutions of light transmittance for typical PMMA/SAN blends filled with SiO2−OH and SiO2−PS particles are shown in Figures 8b and 8c, respectively. For SiO2−OH filled samples, the transmittance decreases a little (by ∼2%) even at low temperatures. The localization of nanoparticles is probably responsible for this decrease. Thus, for the particle-filled blends, we adopted a criterion for the binodal point, 3% reduction of the transmittance, which is a little stronger than the criterion for the neat blends (1% reduction). For PMMA/SAN/SiO2− OH 70/30/3 and 40/60/3 shown in Figure 8b, the binodal temperature with this criterion is estimated to be between 152 and 154 °C and between 156 and 158 °C, respectively. The SiO2−PS particles with PS grafts have higher affinity to PMMA/SAN compared to the bare SiO2−OH particles. However, the particles with PS grafts still exhibit aggregation in the PMMA-rich domain (cf. Figure 2) to reduce the transmittance. For this reason, we adopted the same criterion (3% reduction of transmittance) for the blends filled with those particles to estimate the binodal temperature, 155 and 159 °C for PMMA/SAN/SiO2−PS 70/30/3 and 40/60/3, respectively (cf. Figure 8c). In Figure 6, the binodal phase diagram of PMMA/SAN, PMMA/SAN/SiO2−OH, and PMMA/SAN/SiO2−PS optically determined under the isothermal condition is compared with that determined by rheology. Differing from the rheologically determined diagram, the optically determined diagram (filled symbols) is just moderately affected by the particle, in particular by the bare SiO2−OH particle, possibly because the interaction between the particles and component polymers is not very strong. In relation to the above results, we should keep in mind that the cloud point (binodal point) changes with the criterion to judge the phase separation. In this work, we have adopted the 3% reduction of transmittance as the criterion for the particlefilled systems (a larger reduction compared to the neat blend). Even considering a small uncertainty in this criterion for the particle-filled blends, we can unequivocally conclude that the particles influence the optically determined phase diagram just
Figure 8. Time dependence of transmittance under different temperatures for samples PMMA/SAN 70/30 and 40/60 with (a) no nanoparticles, (b) 3 wt % SiO2−OH, and (c) 3 wt % SiO2−PS.
shows the typical time evolution of transmittance at constant temperatures for PMMA/SAN with or without silica nanoparticles. For near-critical PMMA/SAN 70/30 (left panel of Figure 8a), the transmittance hardly changes during the F
dx.doi.org/10.1021/ma301186b | Macromolecules XXXX, XXX, XXX−XXX
Macromolecules
Article
Figure 9. TEM images of PMMA/SAN without and with silica nanoparticles. Near-critical 70/30 containing (a) no particle (neat blend), (b) 3 wt % of SiO2−OH particles, and (c) 3 wt % of SiO2−PS particles. All these images were obtained after annealing at 172.5 °C for 2 h. Off-critical 40/60 containing (d) no particle (neat blend), (e) 3 wt % of SiO2−OH particles, and (f) 3 wt % of SiO2−PS particles. The images (d) and (e) were obtained after annealing at 180 °C for 2 h and the image (f) after annealing at 181 °C for 2 h. In all images, PMMA-rich and SAN-rich domains are observed as white and dark regions, respectively.
moderately. In contrast, the rheological transition point (apparent binodal point) is strongly affected by the particles and is higher, by a factor >10 °C, than the optically determined cloud point. This difference can be partially related to kinetic aspect of phase separation, as discussed below. 3.4. Phase Growth Kinetics. The rheological transition temperature is much higher than the optically determined cloud temperature in our system (Figure 6), although rheology is believed to be more sensitive than the optical method. One might attribute this difference to a difference in the time scale of the two methods. However, the cloud point hardly changes even if the optical data in Figure 8 at the annealing time of 60 min (∼ time scale of rheological frequency sweep test) are utilized to evaluate this point. Therefore, the difference is to be ascribed to the spatial resolution of the two methods. Resolution of the optical method is determined by the wavelength of light (and also influenced by nanoparticle aggregation, if any). In contrast, resolution of the rheological method strongly depends on relative contributions to the storage modulus from the components and the interface. The interface contribution, utilized as the evidence of phase separation, is clearly observed only when this contribution is comparable to/larger than the component contribution at low frequencies. For blends having large dynamic asymmetry and rather small component moduli, such as PB/LPI blends5 or PS/PVME,4 the interface contribution can be easily resolved so that the rheological method works more sensitively than the optical method. However, this is not the case for our PMMA/SAN blends having just small dynamic asymmetry and large component moduli. For those PMMA/SAN blends, the rheological method serves as a sensitive method only when the phase-separated structure is well developed at high temperatures and after long time annealing. In this sense, the rheological transition temperature depends not only on the component moduli but also on the kinetics of phase growth.
Another question arises for a large difference of the rheological transition temperatures of the neat and particlefilled PMMA/SAN blends (Figure 6). For this question, the phase growth was monitored through TEM observation. The results are discussed below for the near-critical and off-critical blends separately. 3.4.1. Near-Critical Blends. Figures 9a−c show typical TEM images of near-critical PMMA/SAN 70/30 blends after annealing at 172.5 °C for 2 h. Selective locations of bare SiO2−OH and PS-grafted SiO2−PS nanoparticles in PMMA domains and PMMA/SAN interface are clearly seen. The average characteristic length of cocontinuous structure (lc), defined in the previous study,17 was measured during the annealing process. The results are summarized in Figure 10a.
Figure 10. Characteristic length in PMMA/SAN blends with or without silica nanoparticles during the phase separation: (a) nearcritical 70/30 annealed at 172.5 °C; (b) off-critical 40/60 annealed at 180 °C (for neat and SiO2−OH filled blend) and at 181 °C (for SiO2−PS filled blend). G
dx.doi.org/10.1021/ma301186b | Macromolecules XXXX, XXX, XXX−XXX
Macromolecules
Article
The lc value as well as its exponent α (lc ∝ tα) are hardly affected by the particles. This lack of particle effect on lc suggests that the particles hardly change the interfacial contribution to the storage modulus. Thus, the particle effect on the rheological transition temperature can be related to the blend modulus (component moduli in a given morphology). For the near-critical PMMA/SAN 70/30 annealed at 172.5 °C, the modulus at low frequencies is enhanced significantly by the SiO2−PS particles (located at the interface) but just a little by the SiO2−OH particles (segregated in the PMMA rich phase), as noted in Figure 11a. Therefore, for the rheological resolution
domains n in the three blends differs considerably. Specifically, after 8 min annealing, the neat blend has n ∼ 20 droplets/μm2, while the SiO2−PS and SiO2−OH filled blends have n ∼17 and ∼11 droplets/μm2, respectively. The SiO2−OH particles possibly act as nuclei for phase separation owing to the PMMA chains adsorbed thereon, but those particles form rather large aggregates to give a considerably small number of nucleation sites (small n). Correspondingly, there are evident differences in average domain size and its growth rate in the offcritical blends (Figure 10b). In the initial stage, the average domain radius R is in the order of Rneat < RSiO2−PS < RSiO2−OH. But the coarsening exponent β (R ∝ tβ) is in the opposite order (though the difference is small). Relatively large domains would have been formed in the SiO2−OH filled blend because of the nucleation effect of large aggregates of the particles, while the small βSiO2−OH value may be related to the reduced chain mobility. For the SiO2−PS filled blend, Figure 10b also demonstrates that the coarsening suddenly slows (βSiO2−PS decreases from 0.22 to 0.08) in the late stage at t ∼ 30 min. This interesting behavior can be related to the migration of the SiO2−PS particles in the PMMA/SAN interface: For the near-critical SiO2−PS filled blend, the particles were found to be initially dispersed rather randomly in the PMMA-rich phase but migrate to the interface with time (TEM images shown in Figure S.3). This should be the case also for the off-critical SiO2−PS filled blend examined in Figure 10b. Once the particles migrate to the interfaces (surface of the PMMA droplet phase), they would behave as the compatibilizers34 (or stabilizer) to slow the coarsening of the droplet phase. Thus, the sudden slowing seen in Figure 10b can be related to the particle migration to the interface. Of course, the approach of the particles to the interface is rather complex. This approach may be disturbed by the backflow from the interface when the particles come to the vicinity of the interface or influenced by the complicated flow generated during the coarsening. All these hydrodynamic factors may make the real migration more complicated than the ideal diffusion process. However, we may still estimate the migration time of nanoparticles from their diffusion coefficient D0 = kBT/6πηsa, where kB is the Boltzmann constant, T is the absolute temperature, a is the particle radius, and ηs is the matrix viscosity.51 The time required for the nanoparticles to diffuse over their diameter can be estimated as tD = (2a)2/6D0 ∼ 28 min using a = 15 nm and ηs = 240 000 Pa·s (viscosity of PMMA). This diffusion time is close to the time for the onset of slowing, which seems to support the mechanisms of slowing discussed above (stabilization of the droplets due to the particles migrated to the interface). Finally, it should be noted that the storage moduli of offcritical PMMA/SAN 40/60 blend is hardly affected by SiO2− PS particles as well as SiO2−OH particles (Figure 11c). This result suggests that the higher rheological transition temperature seen for the SiO2−OH filled 40/60 blend (Figure 6) is mainly related to the interfacial contribution to the modulus that is affected by the phase growth rate, not to the blend modulus (component moduli in a given morphology) that governs the behavior of near-critical blends. Because the PMMA-rich droplet phases in the SiO2−OH filled 40/60 blend are a little larger but much scarcer than those in the other two blends (Figures 9d−f and 10b), rheological detection of the interfacial contribution requires the phases to grow more significantly in the former blend than in the latter two. The
Figure 11. Storage modulus of PMMA/SAN with and without silica nanoparticles: (a) near-critical 70/30, (b) near critical 80/20, and (c) off-critical 40/60 at temperatures as indicated.
of the interfacial contribution on phase separation, the phases need to grow just to a similar extent in the SiO2−OH filled blend and neat blend but more pronouncedly in the SiO2−PS filled blend. This difference would have resulted in the rheological transition temperature being similar for the first two blends but much higher for the last blend (Figure 6). This explanation can be further tested through comparison of the near-critical PMMA/SAN 80/20 blends filled with SiO2− OH and SiO2−PSl particles. The SiO2−PSl particles having long PS grafts are located in the PMMA-rich phase as similar to the SiO2−OH particles, which enables the comparison without being affected by a difference of the particle location. The modulus at low frequencies for SiO2−PSl filled blend is larger than SiO2−OH filled blend due to different interaction and dispersion of nanoparticles, as noted in Figure 11b. The rheological transition point was found to be in the same order (177.0 °C for SiO2−PSl filled blend > 174.8 °C for SiO2−OH filled blend), which lends support to the above explanation. Thus, the difference of the rheological transition point observed for the near-critical blends can be mainly attributed to the difference of the blend modulus (component moduli in a given morphology) but not to the phase growth rate similar for the three systems (Figure 10a). 3.4.2. Off-Critical Blends. For the off-critical PMMA/PAN 40/60 blend without and with silica particles, Figures 9d−f show droplet-matrix morphology formed via nucleation and growth mechanism. As explained earlier, the SiO2−OH and SiO2−PS particles are selectively located in the PMMA-rich phase and interfaces, respectively. The number density of H
dx.doi.org/10.1021/ma301186b | Macromolecules XXXX, XXX, XXX−XXX
Macromolecules
Article
rheological transition temperature, being higher for the SiO2− OH filled off-critical blend (Figure 6), possibly reflects this difference. Therefore, if the nanoparticles have little effect on the component’s rheology, the change in viscoelasticity during phase growth can be directly correlated with the morphology of blend.
Notes
The authors declare no competing financial interest.
■
ACKNOWLEDGMENTS The authors thank the support from the National Natural Science Foundation of China (No. 50930002 and No. 21074072) and National Basic Research Program of China (973 Program) 2011CB606005. W. Yu is supported by the Program for New Century Excellent Talents in University and the SMC project of Shanghai Jiao Tong University. The authors sincerely thank the anonymous reviewers for their valuable comments and suggestions.
4. CONCLUSIONS The selective location of silica particles was controlled by varying the chain length of polystyrene grafted thereon. The particles with short and long PS grafts were localized at the interface between the component phases and in the PMMArich phase in the PMMA/SAN blends, possibly due to the interfacial tension higher for the PS−SAN pair than for the PS−PMMA pair. Those silica particles hardly affected the optically detected cloud points. However, the widely utilized rheological method gave much higher transition temperatures (apparent binodal temperatures) that changed with the particle type and blend composition. The rheological method detects the interfacial contribution to the modulus, so that this feature of the transition temperature can be related to modulus of the PMMA/SAN blends and the phase growth therein. Specifically, for the PMMA/SAN blends having near-critical compositions (70/30 and 80/20), the cocontinuous phases grew at nearly the same rate irrespective of the type and content of the particles. The blend modulus (component moduli in a given morphology) at low frequencies was enhanced significantly by the PS-grafted particles (located at the interface) but just moderately by the bare particles (having OH groups at the surface and being segregated in the PMMA-rich phase). Thus, for the rheological resolution of the interfacial contribution on phase separation, the phases need to grow to just a similar extent in the blend filled with bare particles and neat blend but more pronouncedly in the blend filled with PS-grafted particles. This difference would have led to the observed rheological transition temperature being similar for the first two blends but much higher for the last blend. In contrast, for the blends with off-critical compositions (≤40/60), growth of the PMMA-rich droplet phases (dispersed in SAN-rich matrix) was affected by the particles. The droplet number density was smaller in the blend filled with bare particles than in the blend filled with PS-grafted particles and/ or neat blend. These blends exhibited similar modulus. Reflecting these difference and similarity, rheological detection of the interfacial contribution on phase separation requires the droplet phases to grow more significantly in the blend filled with bare particles than in the other two blends, which possibly resulted in the observation that the rheological transition temperature was higher for the former blend.
■
■
ASSOCIATED CONTENT
S Supporting Information *
Detailed synthesis procedure and characterization of the modification of SiO2, the analysis of selective location in terms of interfacial tension, and the observation of timedependent migration of silica nanoparticles. This material is available free of charge via the Internet at http://pubs.acs.org.
■
REFERENCES
(1) Larson, R. G. Rheol. Acta 1992, 31 (6), 497−520. (2) Fredrickson, G. H.; Larson, R. G. J. Chem. Phys. 1987, 86 (3), 1553−1560. (3) Niu, Y. H.; Yang, L.; Shimizu, K.; Pathak, J. A.; Wang, H.; Wang, Z. G. J. Phys. Chem. B 2009, 113 (26), 8820−8827. (4) Sharma, J.; Clarke, N. J. Phys. Chem. B 2004, 108 (35), 13220− 13230. (5) Zhang, R. Y.; Cheng, H.; Zhang, C. G.; Sun, T. C.; Dong, X.; Han, C. C. Macromolecules 2008, 41 (18), 6818−6829. (6) El Mabrouk, K.; Bousmina, M. Polymer 2005, 46 (21), 9005− 9014. (7) Lee, H.-N.; Newell, N.; Bai, Z. F.; Lodge, T. P. Macromolecules 2012, 45 (8), 3627−3633. (8) Hobbie, E. K.; Jeon, H. S.; Wang, H.; Kim, H.; Stout, D. J.; Han, C. C. J. Chem. Phys. 2002, 117 (13), 6350−6359. (9) Jeon, H. S.; Nakatani, A. I.; Han, C. C.; Colby, R. H. Macromolecules 2000, 33 (26), 9732−9739. (10) Vlassopoulos, D.; Terakawa, T.; Fuller, G. G. J. Rheol. 2003, 47 (1), 143−161. (11) El-Mabrouk, K.; Belaiche, M.; Bousmina, M. J. Colloid Interface Sci. 2007, 306 (2), 354−367. (12) Chopra, D.; Vlassopoulos, D.; Hatzikiriakos, S. G. J. Rheol. 1998, 42 (5), 1227−1247. (13) Jeon, H. S.; Nakatani, A. I.; Hobbie, E. K.; Han, C. C. Langmuir 2001, 17 (10), 3087−3095. (14) Niu, Y.-H.; Wang, Z.-G. Macromolecules 2006, 39 (12), 4175− 4183. (15) Madbouly, S. A.; Ougizawa, T. Macromol. Chem. Phys. 2004, 205 (9), 1222−1230. (16) Zheng, Q.; Du, M.; Yang, B. B.; Wu, G. Polymer 2001, 42 (13), 5743−5747. (17) Yu, W.; Zhou, W.; Zhou, C. X. Polymer 2010, 51 (9), 2091− 2098. (18) Yu, W.; Bousmina, M.; Grmela, M.; Palierne, J.-F.; Zhou, C. X. J. Rheol. 2002, 46 (6), 1381−1399. (19) Yu, W.; Li, R. M.; Zhou, C. X. Polymer 2011, 52 (12), 2693− 2700. (20) Du, M.; Gong, J. H.; Zheng, Q. Polymer 2004, 45 (19), 6725− 6730. (21) Li, R. M.; Yu, W.; Zhou, C. X. Polym. Bull. 2006, 56 (4), 455− 466. (22) Zou, F. S.; Dong, X.; Liu, W.; Yang, J.; Lin, D. M.; Liang, A. M.; Li, W.; Han, C. C. Macromolecules 2012, 45 (3), 1692−1700. (23) Li, R. M.; Yu, W.; Zhou, C. X. J. Macromol. Sci., Part B 2007, 46 (5), 1051−1062. (24) Li, R. M.; Yu, W.; Zhou, C. X. J. Macromol. Sci., Part B 2006, 45 (5), 889−898. (25) Yu, W.; Zhou, C. Polymer 2012, 53 (3), 881−890. (26) Lee, S. H.; Kontopoulou, M.; Park, C. B. Polymer 2010, 51 (5), 1147−1155. (27) Bose, S.; Ö zdilek, C.; Leys, J.; Seo, J. W.; Wübbenhorst, M.; Vermant, J.; Moldenaers, P. ACS Appl. Mater. Interfaces 2010, 2 (3), 800−807.
AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. I
dx.doi.org/10.1021/ma301186b | Macromolecules XXXX, XXX, XXX−XXX
Macromolecules
Article
(28) Ö zdilek, C.; Bose, S.; Leys, J.; Seo, J. W.; Wübbenhorst, M.; Moldenaers, P. Polymer 2011, 52 (20), 4480−4489. (29) Vandebril, S.; Vermant, J.; Moldenaers, P. Soft Matter 2010, 6 (14), 3353−3362. (30) Vermant, J.; Cioccolo, G.; Golapan Nair, K.; Moldenaers, P. Rheol. Acta 2004, 43 (5), 529−538. (31) Chung, H.-j.; Ohno, K.; Fukuda, T.; Composto, R. J. Nano Lett. 2005, 5 (10), 1878−1882. (32) Chung, H.-j.; Ohno, K.; Fukuda, T.; Composto, R. J. Macromolecules 2006, 40 (2), 384−388. (33) Gam, S.; Corlu, A.; Chung, H. J.; Ohno, K.; Hore, M. J. A.; Composto, R. J. Soft Matter 2011, 7 (16), 7262−7268. (34) Kwon, T.; Kim, T.; Ali, F. b.; Kang, D. J.; Yoo, M.; Bang, J.; Lee, W.; Kim, B. J. Macromolecules 2011, 44 (24), 9852−9862. (35) Walther, A.; Matussek, K.; Müller, A. H. E. ACS Nano 2008, 2 (6), 1167−1178. (36) Vleminckx, G.; Bose, S.; Leys, J.; Vermant, J.; Wübbenhorst, M.; Abdala, A. A.; Macosko, C.; Moldenaers, P. ACS Appl. Mater. Interfaces 2011, 3 (8), 3172−3180. (37) Lipatov, Y. S.; Nesterov, A. E.; Ignatova, T. D.; Nesterov, D. A. Polymer 2002, 43 (3), 875−880. (38) Nesterov, A. E.; Lipatov, Y. S. Polymer 1999, 40 (5), 1347− 1349. (39) Huang, Y. J.; Jiang, S. J.; Li, G. X.; Chen, D. H. Acta Mater. 2005, 53 (19), 5117−5124. (40) Yurekli, K.; Karim, A.; Amis, E. J.; Krishnamoorti, R. Macromolecules 2003, 36 (19), 7256−7267. (41) Yurekli, K.; Karim, A.; Amis, E. J.; Krishnamoorti, R. Macromolecules 2003, 37 (2), 507−515. (42) Ginzburg, V. V. Macromolecules 2005, 38 (6), 2362−2367. (43) Gharachorlou, A.; Goharpey, F. Macromolecules 2008, 41 (9), 3276−3283. (44) Gao, J. P.; Huang, C. W.; Wang, N.; Yu, W.; Zhou, C. X. Polymer 2012, 53 (8), 1772−1782. (45) Chung, H.-J.; Kim, J.; Ohno, K.; Composto, R. J. ACS Macro Lett. 2011, 1 (1), 252−256. (46) Steinmann, S.; Gronski, W.; Friedrich, C. Polymer 2002, 43 (16), 4467−4477. (47) Mo, H. Y.; Zhou, C. X.; Yu, W. J. Non-Newtonian Fluid Mech. 2000, 91 (2−3), 221−232. (48) López-Barrón, C. R.; Macosko, C. W. Soft Matter 2010, 6 (12), 2637−2647. (49) Polios, I. S.; Soliman, M.; Lee, C.; Gido, S. P.; Schmidt-Rohr, K.; Winter, H. H. Macromolecules 1997, 30 (15), 4470−4480. (50) Kasarova, S. N.; Sultanova, N. G.; Ivanov, C. D.; Nikolov, I. D. Opt. Mater. 2007, 29 (11), 1481−1490. (51) Larson, R. G. The Structure and Rheology of Complex Fluids; Oxford University Press: New York, 1999.
J
dx.doi.org/10.1021/ma301186b | Macromolecules XXXX, XXX, XXX−XXX
