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Surface-Influenced Phase Separation in Organic Thin Films on Drying Ravi F. Saraf,*,† Steve Ostrander,‡ and R. M. Feenstra§ T. J. Watson Research Center, IBM Corporation, Yorktown Heights, New York 10598, Department of Material Science, Alfred University, Alfred, New York 14802, and Physics Department, Carnegie Mellon University, Pittsburgh, Pennsylvania 15213 Received October 15, 1996. In Final Form: June 22, 1997 Kinetics of phase separation at the air/polymer interface in a binary polymer mixture on evaporation of common solvent is studied. The lateral dimension of the highly anisotropic, pancake-like, minority phase increases with a growth exponent of 2/3, identical to “late-stage” growth under (classical) thermal quench at interface. The “wetting-layer” formed at the surface is directly visualized using atomic force microscopy. In contrast to the thermal quench, during drying the kinetics depends on the initial condition (i.e., initial concentration, c0) that is resealed to obtain a master curve.
The effects of solvent drying on the morphology of (organic) thin films made up of more than one polymer is of great importance to industries such as microelectronics, paint and coating, and biotechnology (specifically, to make membranes where the pore size is controlled by regulating the phase-separated morphology of the organic film prior to removal of one of the phases). Some of the usual morphological properties of interest are surface smoothness (for low friction or aesthetics), roughness (for high adhesion to the next surface), and film clarity. All these are closely related to the average (lateral) size of the phaseseparated regions, making kinetics during film drying (the usual method for application) an important process for tailoring (organic) thin, multicomponent, film properties. Since the evaporation is at the surface, the phase separation kinetics and the domain growth will be highly influenced by the interracial behavior of the system, in particular the relative wetting/dewetting behavior of the two polymers. To our knowledge there are no reports studying the dynamics of phase separation during drying at the interface. Our results on drying reported here show remarkable quantitative and phenomenological similarity to surface-influenced phase separation on deep thermal quenching (into the unstable region). Recently, there has been a great deal of interest in surface-influenced phase separation (SIPS) phenomena in polymer and monomer systems on thermal quenching.1-6 We will refer SIPS due to thermal quench as classical SIPS to distinguish from our case of SIPS on drying. The characteristic length of one of the phases during separation increases with time as ta, where a is the growth exponent. There are distinctions from the classical 3D (three-dimensional) system where the growth exponent, a ) 1/3, can increase to 1.0 due to coalescence.7 Some general observations from the studies on classical †
IBM Corp. Alfred University. § Carnegie Mellon University. ‡
(1) Guenoun, P.; Beysens, D.; Robert, M. Phys. Rev. Lett. 1990, 65, 2406. (2) Wiltzius, P.; Cumming, A. Phys. Rev. Lett. 1991, 66, 3000. (3) Jones, R. A. L.; Norton, L. J.; Kramer, E. J.; Bates, F. S.; Wiltzius, P. Phys. Rev. Lett. 1991, 66, 1326. (4) Bruder, F.; Brenn, R. Phys. Rev. Lett. 1992, 69, 624. (5) Tanaka, H. Phys. Rev. Lett. 1993, 70, 2770. (6) Tanaka, H. Phys. Rev. Lett. 1993, 70, 53. (7) Siggia, E. D. Phys. Rev. A 1979, 20, 595.
SIPS phenomena are as follows: (i) Due to preferential wetting of one phase over other, there is a low-energy interfacial layer of the latter at the interface as predicted by Cahn.8 The growth kinetics of the characteristic length is also influenced due to the differences in the wetting tendencies of the phases at the interface. (ii) According to Tanaka’s model,5 at an early stage the lateral growth occurs mainly due to two-dimensional (2D) capillary instability with the more-wetting phase transported to the interface leading to a growth exponent of a ) 3/2. This is experimentally observed for a polymeric system where due to slow dynamics the early stage was possible to probe.4 (iii) The late-stage kinetics, observed for small-molecule fluids, indicated a slower exponent of a ) 2/3.1 (iv) Tanaka’s model5 further indicates that the growth exponent for the surface layer thickness should be 1.0 (at least) during the early-stage regime. This seems consistent with the neutron reflectivity results for a binary polymer mixture.3 (v) An interesting observation is recently reported by Cumming et al., where a can range from l to 1.5 depending on the quench depth with respect to the coexistence curve.9 Here we report our studies on phase separation kinetics at the surface in a binary polymer blend as the cosolvent evaporates. Essentially, in the experiment we achieve a “concentration quench”, where the solvent is instantaneously dried (compared to the slow phase separation kinetics of the polymeric system) in the top layer. This layer continues to phase separate until all the solvent from the bulk layer is evaporated. Since the glass transition temperatures (Tgs) of the two polymers chosen are >80 °C above the drying temperature, the system is frozen after solvent evaporation.10 No measurable difference in the lateral size is observed when the sample is left over 60 min beyond the evaporation time, τ.10 By control of the evaporation time, various stages of phase separation are obtained. The characteristic lateral size of the discrete phase in the top layer is measured by atomic force microscopy (AFM). AFM is chosen because, this way we only probe the phase separation in the top layer avoiding the morphology in the bulk. The topography from the deeper layers in the bulk is filtered by a special (8) Chan, J. W. J. Chem. Phys. 1977, 66, 3667. (9) Shi, B. Q.; Harrison, C.; Cumming, A. Phys. Rev. Lett. 1993, 70, 206. (10) Saraf, R. F. Macromolecules 1993, 26, 3623.
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we mention only the salient features. The cosolvent NMP forms a stable complex with PAA and PAETE as shown by points D and E. The corresponding solvent fractions are cD ) 0.1915 and cE ) 0.0946.12,13 The solvent concentration at the onset of phase separation for φ ) 0.2 is cs ) 0.890. D′, E′, and I are given by11
c|x )
Figure 1. Phase diagram of polymers PAA and PAETE and cosolvent NMP at 70 °C. All compositions are in weight percent.
image analysis described below. We first discuss the system chosen to achieve the proper concentration quench. Then we describe the growth curve that is dependent on the initial solvent concentration. Next the observations are discussed in terms of a phenomenological model. Finally, a master curve is obtained from the various growth curves at different initial solvent concentrations. Experimental Section We start with a one-phase, ternary mixture of two polymers, acid (PAA) and ester (PAETE) precursors of poly(pyromellitic dianhydride oxydianiline) (PMDA-ODA) in a cosolvent Nmethylpyrrolidinone (NMP). The weight average molecular weights and polydispersity index for PAA and PAETE are 152 000 and 88 000 and 2.2 and 2.1, respectively. In this study the relative weight fraction of PAA with respect to PAETE, defined as φ, is fixed at 0.2. The weight fraction of NMP in the solution is c0. The c0 ) 0.950, 0.936, and 0.920 for this study. Films from solutions at given c0 are spin coated on a ∼0.2 mm thick single crystal Si wafer with ∼1 nm native oxide. During spin coating at room temperature, an insignificant amount of solvent evaporates (0.0017 cm2/s for Si and polyimide, respectively, the film will attain a hot plate temperature of 70 °C in 10-9 cm2/s. However, the films in these studies are not cured to rigid polyimide structure. Thus, conservatively, at 70 °C, we assume D ∼ 10-9 cm2/s (at most).
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Figure 2. The AFM topograph is obtained on Digital Instrument’s Nanoscope III model at contact force of ∼50 nN range. The 400 × 400 pixel image is reproduced using Galaxy on an RS6000 machine. The range in the x-y plane is 5 µm. The z range corresponding to the highest elevation, (measured) evaporation time, τ, and 〈Dξ〉 for the three samples made from c0 ) 0.916 solutions are as follows: (a) 34.8 nm, 46.0 s, and 676 ( 28 nm; (b) 27.3 nm, 28.3 s, and 473 ( 13 nm; (c) 22.0 nm, 19.5 s, and 354 ( 7 nm.
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Figure 3. AFM micrograph under similar condition as shown for Figure 2. The range in x-y plane is 4 µm. c0 ) 0.934, τ ) 30.24 s, and 〈Dξ〉 ) 418 nm for the film. The nominal height of the mounds is 20 nm.
Results and Discussion The average lateral size, 〈Dξ〉 of the PAA-rich regions is computed by analyzing the 512 × 512 pixel AFM image. From the cross-sectional view, the mounds are shaped close to inverted paraboloids (see Figure 3 and ref 10). Thus, by plotting the curvature of the topography (defined as d2z/dx2 + d2z/dy2 for topography z(x,y)), an image is achieved where the PAA-rich regions are uniformly higher than the background. An example of this is shown in parts a and b of Figure 4 for a film made from c0 ) 0.934 solution with τ ) 26.2 s. A histogram of pixel values for a curvature such as Figure 4b reveals two peaks, one corresponding to the PAA-rich regions and the other to the background mixed phase, as illustrated in Figure 4c. Defining a discriminator level at the midpoint between the two peak positions in the histogram allows a measure of the diameter of the PAA-rich regions. Figure 4d shows pixels with curvature above (below) the discriminator value are shaded white (black). PAA-rich regions which lie on the edge of the image or are noncircular (with ratio of principal moments of inertia outside the range 0.5-2) are excluded and are shown by the gray regions in Figure 4d. Thus both the edge effect and features from phase separation at deeper layer are identified and neglected from the computation. The domains at larger depths from the interface that are observed at the surface as mounds are deciphered as follows: They typically have lower elevation compared to the surface features (see Figure 4b) and appear as either small circular or irregular-shaped features, after the discriminator analysis (see Figure 4d). For example, compare one domain in the lower left corner, one domain in the upper right corner, and one domain in the center, in images b and d of Figure 4. In Figure 4b these domains appear circular but with lower elevation. The corresponding features in Figure 4d are noncircular and thus colored gray, i.e., dropped from the analysis. To also note is that there are similar feature in Figure 4b adjacent to the domains in the lower left and upper right corner that are not dropped in Figure 4d. However, they are small and do not affect the 〈Dξ〉 computation (discussed in the next paragraph). Finally, the Dξ of the PAA-rich regions is obtained by the area A of the white regions in Figure 4d, using Dξ ) 2(A/π)1/2. Figure 4e shows the distribution curve corresponding to Figure 4d.
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The typical size distributions seen in Figure 4e (are likely to) occur due to time lags in the initiation of the unstable fluctuations. Although the quench depth in surface layer is fairly uniform (as discussed in the next paragraph), the small concentration gradient in solvent concentration will cause phase separation at greater depths to initiate later leading to a distribution of lateral size, Dξ, of the PAA-rich mounds. This implies that the larger domains have better correspondence with the total time available, τ, for phase separation. Thus, the characteristic lateral size scale, 〈Dξ〉, for a given τ, is computed by averaging on the large domains. Specifically, the images are split into four quadrants, and the largest size PAA-rich region in each quadrant is found. The mean 〈Dξ〉 and standard deviation of these largest regions are computed, and these are the values plotted in Figure 5. The shape of the growth curve (specially the growth exponent) will be sensitive to the computed 〈Dξ〉 from the distribution curve. We note that although the distribution (Figure 4e) has a broad tail at low sizes, it has a fairly sharp cutoff close to the maximum Dξ population. Thus, the computation of 〈Dξ〉 will be close to the cutoff with small standard deviation. This is apparent from the error bars in Figure 5. Furthermore, the larger domains that initiated first are closer to the surface. Therefore the growth curve observed has better correspondence to the (near) surface behavior. We note in passing that due to interference from the bulk layers, edge effects, large number densities, and complex size distribution, a simple Fourier analysis method is inadequate to quantify 〈Dξ〉 as a function of τ.11 In Figure 5, the growth curve has three distinct regimes: (i) the induction period, where the dried film looks smooth under AFM with no visible sign of phase separation; (ii) the early stage, where it is difficult to probe because of limitation in deciphering the PAA-rich mounds from the film roughness; (iii) the late stage, which shows a distinct growth exponent of 2/3 independent of initial condition, c0. Another remarkable observation is that, for a given τ, film with larger c0 has smaller 〈Dξ〉! This seems counterintuitive since higher c0 should lead to more mobility implying larger 〈Dξ〉. The theoretical framework is very limited for SIPS during thermal quenching, and not present (to our knowledge) for separation during drying. We surmise the following phenomenological model (based on classical SIPS studies5) to explain the observed growth behavior in Figure 5. Let γA, γB, and γAB be surface energies of PAA/air, PAETE/air, and PAA/PAETE interfaces. Since PAA is a strong acid with one acid group per monomer unit and B is its ester form, γA > γB. Since B does not spread on A film, γAB > γA - γB. Initially, as the solvent dries in DTL, the PAETE will begin to form a low-energy surface layer (LESL) (similar to classical SIPS13). The growth of the phase would be by flow through bicontinuous, percolating conduits connecting the bulk to the surface, as suggested and observed by Tanaka for classical SIPS.5 Thus, during the induction period this initial LESL formation process occurs with no visible phase separation in the in-plane direction. Since our system is PAETE rich, the coverage of the LESL is 100%. However, LESL can directly be observed by making a film from a φ ) 1.5 mixture where incomplete coverage of LESL is evident from the circular perforations observed as “holes” in Figure 6. The estimated average thickness of LESL is ∼2.5 nm in Figure 6. Since LESL thickness is an order of magnitude smaller than l, the DTL can be treated as bulk during the induction period.
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Figure 4. An image analysis consisting of five steps are performed to obtain the average size, 〈Dξ〉 of the PAA-rich phase. (a) A (raw) 10 × 10 µm AFM image, the gray level corresponds to height (i.e., topography). (b) Image is replotted with pixels corresponding to curvature, defined in the text. (c) Curvature distribution of the pixels. For a discriminator curvature value of 152 (defined in text), image d is obtain. The pixels above and below 152 are shaded white and black, respectively. Part e shows the size distribution of the discrete white (PAA-rich) regions in (d) corresponding to, 〈Dξ〉 ) 343 ( 5 nm at τ ) 26.2 s.
Similar to the classical SIPS,1,5 accompanying the PAETE-rich (E′ phase) LESL a thinner PAA-rich (D′
phase) lower layer in its contact is also formed during the induction period. From Figure 1, since the D′ phase with
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Figure 5. Growth curve, 〈Dξ〉, versus τ showing three regimes: an induction period below which no mounds of PAA-rich phase are observed; an early-stage growth period where fast growth rate seems to occur; the late-stage growth period where 〈D〉 ∼ τ2/3. The various regions are defined approximately and the 2/3 slope curve is drawn simply to guide the eye. The growth curves are for films made from two initial concentrations, viz., c0 ) 0.95 and 0.936.
Figure 6. (a) A 4 × 4 µm AFM micrograph of a 60/40 PAA/ PAETE blend film made form solution with c0 ) 0.940. (b) The cross-sectional view indicated is a vertical line slightly on the left side from the center. The depth of the “holes” is ∼2.5 nm, from Figure 4b.
φ ) 0.85 has a significantly larger amount of PAA than the E′ phase with φ ) 0.03, we expect that γD′E′ > γD′ -
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Figure 7. Master growth curve made from solutions of c0 ) 0.916, 0.934, and 0.950. The vertical shift and horizontal shifts of the other two curves made relative to the curve for c0 ) 0.950 are discussed in the text. The lines are just to guide the eyes with respect to expected scaling laws.
γE′, where, γD′, γE′, and γD′E′ are respective surface energies for D′ and E′ phases. As the PAA-rich phase thickens, eventually, the system will minimize the D′-E′ contact by breaking the D′ phase, to reduce the free-energy. Analogous to classical SIPS,5 the two phases will continue to grow rapidly by transport through the conduits and the (capillary) in stability (i.e., break up of surface layer), in the early-stage period. Figure 5 is qualitatively consistent with this scenario. However the present experimental set-up limits us to probe this early-stage regime in more detail. After the bicontinuous conduit morphology breaks down the coarsening of the discrete phase will slow down, leading to a late-stage regime.1,5 This late-stage coarsening of the discrete D′ phase is distinctly seen in the growth curve (see Figure 5) as a reasonably sharp change in the growth rate. The growth law in the late stage can be written as, 〈Dξ〉 ) K(t/κ)2/3. Unlike classical case, where the scaling factors, K and κ, for correspondence condition (and universality) depend only on the temperature difference between the final (in the unstable region) and the phase separation on-set temperatures, in the drying case they depend on initial condition (i.e., c0) also. The characteristic time, κ is proportional to the dynamics of transport in the conduit, i.e., flow rate, Q ∼ γD′E′/η.5 As c0 increases η increases causing concomitant decrease in κ. We assume, η ∼ c0 leading to κ-1 ∼ c0. The front-factor, K, is related to change in the amount of material in DTL. The longer the time available for DTL to form (i.e., ti), the larger will be the K. We assume K ∼ ti. If li is the hypothetical thickness of DTL before accounting for solvent shrinkage, then ti ∼ li2 or K ∼ li2. Since the amount of polymer available to phase separate in l1 is ∼li(1 - c0), K ∼ l12(1 - c0)2. As discussed earlier, li is nominally constant (∼1 µm), leading to K ∼ (1 - c0)2. Using c0 ) 0.950 as a reference, we rescale the growth curves for c0 ) 0.934 and 0.916 by applying the above mentioned relationships for K and κ. The two curves are normalized by dividing 〈Dξ〉 by (0.066/0.05)2 and (0.084/ 0.05)2, and dividing time, τ by 1.2 and 1.6, respectively, for c0 ) 0.934 and 0.916. The increase in κ with decrease in c0 is consistent with the above discussion. The relative change in κ was calculated empirically by fitting the horizontal shift. Figure 7 shows the resealed curve. Note that the film with larger c0 having smaller Dξ (seen in Figure 5) is rescaled consistently to obtain the master
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curve. The visual inspection of data clearly indicates that the growth exponent is invariant and the rescaling is reasonably good. The good correspondence law to rescale the curves made from different c0 and excellent agreement with the late-stage growth exponent for classical SIPS1 suggest that the 2/3 growth exponent for SIPS due to drying is a universal behavior. We note that although the time range is small (limited by the thickness discussed earlier), the slope is distinctly different from 1/3, 1, and 3/2 exponent lines drawn to aid the eye. Note that at longer time, the growth data in Figure 7 for film made from c0 ) 0.916 follows the 1/3 power law line. We attribute this to large thickness of the films: In thicker films, the phase separation from the bulk superimposes on the surface features. As a result, for longer time (τ) and larger thicknesses, the topography is dominated by the bulk rather than the surface phenomena. This changes the growth exponent to classical 1/3 law. Since the coarsening is faster than classical diffusionlimited processes (i.e., Ostwald ripening16), it is expected that the growth process is still by coalescence process. The coalescence is apparent in Figure 2 for all three τ values shown. However, coalescence of the nonwetting phase (i.e., D′) in the SIPS process is limited in the film (16) de Gennes, P. G. J. Chem. Phys. 1980, 72, 4756.
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plane (i.e., 2D) evidenced by the anisotropic shape. This qualitatively explains the lower growth exponent of 2/3 in contrast to 1.0 observed for the classical 3D coalescence process.7 In summary, we have measured the kinetics of phase separation at the air/film interface during drying. The time sequences of phase separation were obtained by controlling the evaporation time, τ, and allowing the phase separation to occur only during this drying time. Taking advantage of slow dynamics in polymeric fluids, an ideal concentration quench was ensured by achieving deep concentration quench relative to the concentration variation in DTL formed by solvent evaporation within an induction period where no lateral phase separation is observed. Three distinct kinetic regimes are observed in the growth curve: an induction period, an early stage, and a late stage. In the late stage, the size of the nonwetting phase, 〈Dξ〉 ∼ τ2/3. The growth curves at different initial solvent concentrations, c0, can be resealed to superimpose forming a master curve. The consistent resealing parameters to obtain the master curve and identical late stage growth exponent as the classical SIPS process suggest that the observed behavior may be universal. LA960987B
