INTERFACE SCIENCE 11, 225–235, 2003 c 2003 Kluwer Academic Publishers. Manufactured in The Netherlands.
Hierarchic Structure Formation in Binary and Ternary Polymer Blends MONIKA SPRENGER∗ AND STEFAN WALHEIM† Fakult¨at f¨ur Physik, Universit¨at Konstanz, D-78457 Konstanz, Germany ANDRZEJ BUDKOWSKI Smoluchowski Institute of Physics, Jagiellonian University, ul. Reymonta 4, 30-059 Krak´ow, Poland ULLRICH STEINER Department of Polymer Chemistry and Materials Science Center, University of Groningen, Nijenborgh 4, NL-9747 AG Groningen, The Netherlands
[email protected]
Abstract. The phase morphology of multi-component polymer blends is governed by the interfacial interactions of its components. We discuss here the domain morphology in thin films of model binary and ternary polymer blends containing polystyrene, poly(methyl metacrylate), and poly(2-vinylpyridine) (PS, PMMA, PVP). When sandwiched between two non-polar surfaces, characteristic lateral phase morphologies are observed after the film formation by spin-coating. We discuss here two techniques, by which hierarchical lateral structures in polymer films can be made. The first method makes use of two simultaneously occurring interfacial instabilities. The second technique employs the effect of a variation of the enthalpic interaction parameters in a ternary polymer mixture on its lateral polymer phase morphology. Keywords: polymer blends, phase separation, polymer films, atomic force microscopy 1.
Introduction
Commercial polymer based materials are usually complex. They consist of a large number of components, which can be other polymers, low molecular weight organic additives, or non-organic fillers. The large number of components are dictated by the desired bulk properties, such as mechanical toughness, thermal expansion, but also by their surface properties, such as adhesion, abrasion, optical appearance, etc. While multi-component mixtures are routinely addressed by the engineering literature, there are fewer studies that consider such systems from a fundamental point of ∗ Present address:
Max-Plank Institut f¨ur Metallforschung, Heisenbergstrasse 1, D-70569 Stuttgart, Germany. † Present address: Institut f¨ ur Nanotechnologie, Forschungszentrum Karlsruhe, Postfach 3640, D-76021 Karlsruhe, Germany.
view. This is partially due to the rapid increase in complexity for each added component, but also because multi-component systems are often considered as an extension of the much studied binary case. As a model system, we considere here mixtures of mutually incompatible polymers [1, 2]. While the interplay of demixing and wetting in thin film of weakly incompatible polymer blends is reasonably well understood [3], the situation is much less clear in the more common case of strongly incompatible polymers. In most experimental studies, a polymer blend is prepared in the mixed state and demixing is initiated by a temperature quench. Since the viscosities of the polymers are high, the change in phase morphology is quasi-static. It is therefore possible to assign a Gibb’s free energy to the blend at each stage during the phase separation process. In the case of strongly incompatible polymers, this approach is no longer possible. Since the critical
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point of mixtures of most generic polymers lies above their degradation temperature, it is not possible to prepare an initial mixed state. Instead, demixing occurs during the film preparation process, which typically involves the extraction of a solvent (e.g. in a spin-coating experiment). While in both cases the change in phase morphology is primarily driven by the various interfacial energies in the system, the situation is complex for phase separation during a quench by solvent extraction. All parameters that are inherent to the experimental system (e.g. the surface and interfacial tension, viscosities) change in a continuous, but unknown fashion. Since the diffusion kinetics of the solvent vapor into the atmosphere is not known, it is difficult (if not impossible) to describe phase separation during demixing in terms of a free-energy model. In addition, film formation by spin-coating involves several hydrodynamic regimes and diffusive processes: an initial low viscosity stage during which the build-up convection rolls is possible [4, 5], an intermediate stage, during which material is rearranged by both convective and diffusive processes, and a high concentration stage during which diffusion is suppressed due to the high incompatibility of the polymer couple [1]. Finally, since the mutual solubility of two strongly incompatible polymers is close to zero, film preparation by spin-coating may lead to long-lived metastable phase morphologies [1]. Therefore, despite its technological importance, a quantitative model for the demixing of strongly incompatible polymer blends is not available and we have to resort to a qualitative description instead. In spite of these complications, the characteristic phase morphology of a binary blend is determined (in a quantitatively unknown fashion) by the mutual enthalpic interaction parameters of the blend components. If three liquid phases are involved, the arrangement of the three phases can be described in terms of a wetting argument, in which one of the components completely or partially wets the interface of the two most incompatible components [6–8]. By altering the balance of the three liquid-liquid interfacial energies, the morphology of the intercalated liquid should change. In particular, when going through a wetting transition, a significant change in the arrangement of the three phases is expected. Here, we describe an experimental study of the phase morphologies of two and three component polymer blends as a function of their mutual incompatibility or, equivalently, as a function of their interfacial energies.
To more conveniently observe the system, we monitor the phase morphology in a thin film, rather than performing bulk measurements. This may introduce complications because of a possible preferential adsorption of one of the components at the surfaces of the film. In previous studies, [1, 2] it was demonstrated, however, that that this effect can be minimized by rendering the substrate surface non-polar. The work reported here is an extension of earlier studies [1, 2, 9, 10], where we have discussed the phase morphology of polystyrene/poly(methyl metacrylate) (PS/PMMA) and polystyrene/poly(methyl metacrylate)/poly(2-vinylpyridine) (PS/PMMA/PVP) blends. Here, we discuss the interplay of two simultaneously occurring interfacial instabilities during the spincoating of binary blends and we explore the effect of modifying the interfacial interactions in a ternary mixture by the incorporation of a low molecular weight solvent into the blend.
2.
Experimental Methods
2.1.
Materials
Polystyrene (PS), poly(methyl metacrylate) (PMMA), and poly(2-vinylpyridine) (PVP) were purchased from Polymer Standards Service, Mainz, and used as obtained (Fig. 1). The molecular characteristics of the polymers are summarized in Table 1. The solvents used were analytical grade tetrahydrofuran (THF), stabilized with 250 mg/l 2.6-di-tert-butyl-4-methylphenol) and Table 1. Polymer PS
Molecular parameters of the polymers. Molecular weight (kg/mol)
Polydispersity Mw /Mn
94.9
1.06
PMMA
126
1.04
PVP
115
1.03
Figure 1. Molecular formulae of the polymers. left: polystyrene (PS), middle: poly(methyl metacrylate) (PMMA), right: poly(2vinylpyridine) (PVP).
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Figure 2. Schematic representation of the film formation during spin-coating. Starting from a homogeneous mixture (a), phase separation sets in as soon as a certain critical polymer concentration is surpassed (b). At this stage all three phases are solvent swollen and the film surface is essentially flat. Due to the differing solubility of the polymers, the solvent concentration in the different phases varies. As solvent evaporation continues, one of the polymers solidifies first, followed by a collapse of the other two polymers (c). This leads to a film topography, which mirrors the lateral arrangement of the polymer phases. As opposed to an equilibrium morphology (obtained e.g. after annealing the sample) the obtained structure does not necessarily minimize the polymer-air surface free energy.
analytical grade cyclohexanone. The polymers were dissolved in one of the solvents (1–3 weight percent polymer), and films were made by spin-coating. The film thickness was controlled by varying the spincoating speed (3000–7000 rpm). Film thicknesses were typically ∼100 nm. 2.2.
Substrate Preparation
The substrates used were highly polished silicon wafers. As obtained, the silicon wafers were covered by a thin (1–2 nm thick) oxide layer. The substrates were cleaned by the snow-jet method [11]: the wafers were exposed to a jet of CO2 ice crystals, which were produced by expanding CO2 through a nozzle. This way, surface contaminants are removed either by mechanical impact, or by dissolution in the CO2 . The clean substrates were first covered by a 2 nm thick chrome layer, followed by ∼30 nm of gold (Au). Since Au surfaces contaminate rapidly under ambient conditions, they were used immediately after evaporation. To produce non-polar surfaces, the Au substrates were immersed over night in a 0.285% solution of octadecylmercaptan in an ethanol-THF mixture (5:2). This results in the formation of a physisorbed self-assembled monolayer on Au. The quality of the resulting substrates was monitored by measuring the contact angle of a water drop that was placed onto the surface. Only substrates with water contact angles larger than 105◦ were used. 2.3.
distribution could be obtained in three different, complementary ways: 2.3.1. Topography Analysis. AFM topography images of the as cast films were not flat, but exhibited a distinct topographic contrast that could be identified with the lateral arrangement of the PS, PMMA and PVP phases. The origin of this contrast (as described earlier [1, 2]) is explained in terms of the relative solubility of the three polymers in the spin-casting solvent (Fig. 2). While THF and cyclohexanone are solvents for PS, PMMA, and PVP, their relative solubility differs. For both solvents, PMMA is the least soluble polymer. During spin-coating, the solvent starts to evaporate and once a critical polymer concentration is surpassed, phase separation sets in. While at this stage there is solvent in all three phases, the solvent concentration differs according to the relative solubility in the polymer. The solvent concentration is higher in the PSand PVP-rich phases and lower in the phase rich in PMMA. As solvent continues to evaporate, a point in time is reached, when the PMMA is solvent free and has therefore solidified in its glassy state, while the other two polymers still contain solvent. As evaporation continues, the PS- and PVP-rich phases collapse below the height of the PMMA phase, forming a hierarchy in local film thickness. The process is similar for the two solvents. By knowing (or experimentally characterizing) the solubilities of the polymers, it is therefore straight forward to interpret the AFM topography images in terms of the local composition of the film.
Atomic Force Microscopy
The ternary polymer films were analyzed using a self-built atomic force microscope (AFM), operated in the contact mode. Information on the lateral phase
2.3.2. Friction Contrast. The friction of an AFM tip when moved across a surface (in contact mode AFM) causes torsional bending of the AFM cantilever. The amount of torsional bending is therefore a measure for
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the friction coefficient of the AFM tip on the local sample area, allowing to identify the lateral distribution of the three polymers. PVP exhibits a significantly higher friction coefficient compared to PS and PMMA. This is mainly due to the hygroscopic nature of PVP: the PVP phase swells slightly in ambient air of sufficiently high humidity. This causes a softening of the PVP phase and therefore a high polymer-tip friction coefficient. The friction coefficients of PS and PMMA are comparable, with the PS friction coefficient slightly higher compared to PMMA. Similar to the topographic information, the friction contrast reveals the lateral composition of the polymer film [9]. 2.3.3. Selective Dissolution. The disadvantage of AFM and related techniques is that only the composition of the sample surface can be probed. In thin films, this information can be supplemented by selectively removing some of the polymers [1, 2, 12]. To this end, the sample is first imaged by AFM. It is then exposed to a selective solvent that dissolves only one of the polymers. In our case the sample was washed for several minutes in ethanol, which is a solvent for PVP, but not for PS and PMMA. Subsequently the sample is dried and an AFM image is taken at the same location as the previous scan. This procedure is repeated a second time with a selective solvent for one of the two remaining polymers (e.g. cyclohexane, which is a selective solvent for PS). The composition information of the entire film is then obtained by superposing the three AFM scans. This type of data analysis confirms the quasi two-dimensional nature of our experimental system: the lateral phase distribution imaged at the film surface extends through the film to the substrate surface. In particular, there is no measurable enrichment (or specific adsorption) of any of the polymer phases at the two film surfaces. 3. 3.1.
Results and Discussion Structure Formation in Binary Polymer Blends
While the structure of binary and ternary polymer mixtures was previously analyzed in terms of equilibrium thermodynamic arguments [1, 2], little was said in terms of the structure formation process. Information on the lateral structure formation process is hard to obtain, since a direct observation of the film formation during spin coating is difficult. We instead try to infer information from the morphology of the final
films and deduce the film formation processes from the morphologies that are “frozen-in” for various film formation parameters. We start by comparing phase morphologies for for all possible binary combinations of PS, PMMA and PVP. In Fig. 3, films of PS/PMMA (a), PMMA/PVP (b) and PS/PVP (c) mixtures ares shown. Apart from the differing polymer ingredients, all other parameters were kept constant. All films were cast onto SAM covered Au substrates from a THF solution. The polymer mixing ratio was 1:1 by weight. The examination of the three images (Fig. 3(a)–(c)) reveals two characteristic length scales in the sample topography. While present in all three images, the relative importance of these length scales varies. The PS/PMMA mixture in Fig. 3(a) exhibits predominately a finely dispersed phase morphology showing domain sizes on the order of λ1 ∼ 1 µm. Careful examination of the image shows a long wavelength undulation (λ2 ∼ 40 µm) superposed on the fine PS-PMMA phase morphology. While this is only a small effect in the case of Fig. 3(a), the surface topography of a PMMA/PVP film in Fig. 3(b) is dominated by the long wavelength undulation that is superposed on the topographic pattern stemming from the PMMA/PVP phase morphology. In Fig. 3(c), the demixing of PS and PVP has formed a hierarchical pattern, in which the both length scales are equally important. It is instructive to compare the three images (Fig. 3(a)–(c)) in terms of the two length scales. While the long wavelength undulation is influenced only little by the nature of the polymers in the film (λ2 lies in the range 35–40 µm), the small length scale λ1 is strongly influenced by the polymer polymer interaction paramter χ . It is smallest for for the PMMA/PVP couple and largest for PS/PVP. Since the polymerpolymer incompatibility increases from PMMA/PVP (χ ≈ 0.007) via PS/PMMA (χ ≈ 0.02) to PS/PVP (χ ≈ 0.1) [2], λ1 seems to scale with with the polymer compatibility. From Fig. 3(a)–(c), it is evident that two competing instabilities govern the structure formation process: a short wavelength instability (λ1 ) that varies with the polymer-polymer interactions and a larger length scale (λ2 ) that seems to be intrinsic to the film formation process. Commonly, λ1 is identified with a (bulk) phase separation process, which is confined into two dimensions by the film formation during spin coating. Alternatively, phase separation could lead initially to a layered arrangement of the two phases, which later
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Figure 3. PS/PMMA (a), PMMA/PVP (b), PS/PVP (c), and PS/PMMA/PVP (d) films spin-cast from THF onto SAM covered Au substrates. The average film thickness was ∼200 nm. The contrast in the AFM images corresponds to a height variation of approx. ±60 nm. The binary polymer films show two charateristic length scales (λ1 and λ2 ), which are expressed to a varying degree. The ternary film in (d) shows only the λ1 instability. The lateral phase morphology shows coexisting PS and PVP phases, with the PMMA intercalated at the PS/PMMA interface.
breaks up via an interfacial instability. While the detailed mechanism is difficult to elucidate based on the images in Fig. 3 alone, more extensive work on binary thin polymer films supports the second scenario [13, 14]. The origin of λ2 is more difficult to determine. Experimentally, this type of surface instability is typically observed after spin-coating from highly volatile solvents (e.g. acetone, chloroform). While other origins of this instability (such as dewetting instabilities or polymer-solvent demixing) cannot be excluded from the present data, it is likely that the λ2 instability is caused by a hydrodynamic effect that is triggered by the rapid solvent evaporation. Extremely high evaporation rates cause a significant cooling of the free surface, possibly causing a Marangoni-B´enard convection instability [4, 5]. In an alternative model, de Gennes [15] proposed recently that the build-up of a polymer concentration gradient across the film could lead to a
convection instability driven by a surface tension difference that arises from this concentration gradient. The schematic drawing in Fig. 4 illustrates the interplay of two interfacial instabilities. Initially, phase separation leads to a bilayer (or possibly as trilayer) configuration. As the solvent evaporates all surface-and interfacial tensions change continuously. This may lead to a situation, in which the bilayer configuration is no longer stable. The decay of the liquid-liquid bilayer can proceed via two interfacial instabilities: a undulatory instability of the liquid-liquid (or polymer-polymer) interface or an instability of the film surface. Since the film surface tension is larger than the interfacial tension between the two phases, we expect the surface instability to exhibit a larger wavelength compared to the interfacial instability. Also, since the surface tension of the film is influenced only little by the composition of the film at the surface, there should be only a small
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Figure 4. Schematic representation of possible scenarios of the structure formation process during spin-coating. Initially, phase separation creates a double layer of the two (highly diluted) phases. As solvent continues to evaporate, two instabilities can develop. An instability of the liquid-liquid interface, caused by the increasing unfavorable enthalpic interactions as the solvent concentration decreases, leads to the break up of the film. Alternatively, an instability of the free surface (caused for example by a hydrodynamic instability, which leads to a lateral variation in polymer composition) leads also to a laterally heterogeneous phase morphology. Both instabilities can take place simultaneously. We tentatively identify λ1 with the interfacial instability (left) and λ2 with the instability of the free surface (right).
variation of the instability with the chemical nature of the polymers in the film. This is, however, different for the liquid-liquid interfacial instability. Since the interfacial tension is a function of the interaction parameter χ , the wavelength of the corresponding instability should vary strongly with polymeric system. We therefore identify λ1 with the instability of the liquid-liquid interface and λ2 with the instability of the film surface. Both wavelengths should vary with layer thicknesses in a characteristic fashion. The film thickness dependence of capillary instabilities depends on the details of the destabilizing force [16] and is usually a strong function of the layer thickness [1]. Convection instabilities, on the other hand, reflect the film thickness at the time of the onset of convection rolls, which is larger by two orders of magnitude [4] compared to the thickness of the final films. The different film topographies in Fig. 3(a)–(c) are a result of the interplay of the two instabilities as illustrated in Fig. 4. For the PS/PMMA and PMMA/PVP mixtures, λ1 is dominating, leading to the charac-
teristic lateral phase morphology. The superposed λ2 instability is visible but has not lead to a film breakup. The stronger relative role of λ2 in the case of the PMMA/PVP mixture (as compared to PS/PMMA) may lie in a different rheology (i.e. a different viscosity) of the two composite films. In the case of the PS/PVP system (Fig. 3(c)), both instabilities play an equally important role during the film formation, leading to a hierarchical structure. The interplay of λ1 and λ2 depends sensitively one the details of film formation. In Fig. 5(a)–(c), PS/PMMA, PMMA/PVP and PS/PVP were cast from a more dilute THF solution (2% by weight) at a higher spin-coating speed (7000 rpm). This resulted in a more rapid film formation and in thinner films, with an average film thickness of 70 nm (as opposed to 200 nm in Fig. 3). Thinner films (corresponding to a more rapid film formation process) result in a reduction of λ1 (as observed by us earlier [1]), and in a complete suppression of λ2 . A more rapid film formation corresponds to a more rapid (and deeper) quench by solvent
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Figure 5. PS/PMMA (a), PMMA/PVP (b), and PS/PVP (c) films spin-cast from THF onto SAM covered Au substrates (2% polymer by weight, 7000 rpm). The average film thickness was ∼70 nm. The contrast in the AFM images corresponds to a height variation of approx. ±15 nm. All three films exhibit a phase morphology characterized by the length λ1 only. The formation of surface waves with the wavelength λ2 is suppressed (as confirmed by AFM scans with a larger image area), presumably by the more rapid film formation process.
evaporation, resulting in a larger chemical potential difference between the mixed and the two-phase configuration when demixing sets in, which in turn should lead to a smaller wavelength of the liquidliquid interfacial instability. Apparently, the faster film formation suppresses λ2 , possibly because film formation is terminated before the instability of the free film surface can develop. Clearly, by appropriately adjusting the spin-coating parameters, it is possible to select one of the two intrinsic length scales. Alternatively, films that exhibit both length scales to a varying degree can be made this way. 3.2.
Structure Formation in Ternary Polymer Blends
The phase morphology of the ternary PS/PMMA/PVP mixture in Fig. 3(d) closely resembles the phase morphologies of the three binary mixtures. In Fig. 3(d)
the long wavelength instability (λ2 ) is completely suppressed. The value of λ1 ≈ 3.5 µm lies between the PS/PMMA phase morphology (λ1 = 1–2 µm) and the more coarse PS/PVP phase morphology (λ1 ≈ 6 µm). This is due to the interfacial activity of PMMA in the three component mixture, which “emulsifies” the PS/PVP mixture, leading to a significantly reduced length scale λ1 compared to the binary PS/PVP mixture in Fig. 3(c). By intercalating at the PS/PVP interface, the PMMA phase eliminates the enthalpically unfavorable PS/PVP interfaces. While this is not easily seen in Fig. 3(d), since the gray levels corresponding to PMMA and PVP are similar, this is easily verified by selectively dissolving some of the polymers [2]. 3.2.1. Wetting at the Polymer-Polymer Interface. The requirement for the interfacial activity in
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Fig. 3(d) is the complete wetting of the PS/PVP interface by PMMA [6–8]. This implies that the sum of the interfacial energies per unit area (i.e. the interfacial tensions) of the PS/PMMA and the
PMMA/PVP interfaces must be lower compared to the PS/PVP interface: γ12 + γ23 < γ13
(1)
Figure 6. PS/PMMA/PVP films (in a 3:1:1 mixing ratio) spin-cast from cyclohexanone onto a SAM covered Au surface (3% polymer, 3000 rpm) at ambient conditions (≈60% humidity). The average film thickness was ∼100 nm. The contrast in the AFM topography images corresponds to a height variation of approx. ±50 nm. (a) topographic image (b) friction image, as cast. In (b), the dark domains correspond to a high friction coefficient. In (c) the PVP phase was removed by dissolution in ethanol, followed by the dissolution of PS in cyclohexane (d), revealing PMMA rings on the substrate. In (e) and (f), higher magnifications images are shown. The topography image in (e) has an enhanced contrast compared to (a) and the friction image in (f) has an inverted friction contrast compared to (b), with bright regions corresponding to a high friction coefficient.
Hierarchic Structure Formation
where the indices refer to PS, PMMA and PVP, respectively. The interfacial energy between two polymer phases γi j is related to the interaction parameter χi j [17] χi j γi j = k B Tρa (2) 6 with a the statistical segment length and ρ −1 the monomer unit volume. Since a and ρ, are comparable for the three polymers, Young’s equation (Eq. (1)) can be recast in terms of the three χ -parameters. By changing one (or several) of the χ -parameters, so that γ12 + γ23 > γ13
(3)
the condition for complete interfacial wetting is no longer fulfilled and a phase morphology different from
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Fig. 3 is expected. Since PS, PMMA, and PVP are mutually incompatible (and since therefore the three χ -parameters cannot be sufficiently modified by varying the temperature), changing the polymer-polymer interaction parameter implies changing the chemical nature of one of the phases. This principle is illustrated in Figs. 6 and 7, where a PS/PMMA/PVP mixture (3:1:1) was spin-cast onto a SAM covered Au surface from a cyclohexanone solution (3% polymer by weight, at 3000 rpm). While cast from the same solution onto a similar substrate, the only difference in the preparation conditions was the relative humidity, at which the films were prepared. The films in Fig. 6 were made at ambient conditions (room temperature (≈22◦ C), relative humidity ≈60%). For the preparation of the samples in Fig. 7, the humidity in the spin-coater was lowered to less than 20% by covering the spin-coating
Figure 7. Identical sample preparation as in Fig. 6, except for a lowered humidity (
