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Mar 21, 2016 - Effects of Polymer Tacticity and Molecular Weight on the Glass. Transition Temperature of Poly(methyl methacrylate) Films on Silica. Ku...
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Effects of Polymer Tacticity and Molecular Weight on the Glass Transition Temperature of Poly(methyl methacrylate) Films on Silica Kun Geng† and Ophelia K. C. Tsui*,†,‡ †

Department of Physics, Boston University, Boston, Massachusetts 02215, United States Division of Materials Science and Engineering, Boston University, Boston, Massachusetts 02215, United States

‡

S Supporting Information *

ABSTRACT: The glass transition temperature (Tg) of poly(methyl methacrylate) (PMMA) films supported by silica was studied as a function of film thickness at two molecular weights (MW = 2.5 and ∼50 kg/mol) and two syndiotacticities (s% = 56 and 69 or 79%). We found that the result depends on both the MW and s% in unconventional ways. For the 2.5 kg/mol films, the Tg was depressed for the lower s% films but enlarged for the higher s% films. For the ∼50 kg/mol films, however, the signs of the Tg change were reversed. To understand these observations, we constructed bilayer films from pairs of the studied PMMAs with roughly the same MW but different s%. Our result suggests that for the higher MW films, both the effects of the free surface and the substrate interface on the Tg are more significant on the higher s% PMMAs than on the lower s% PMMAs. For the lower MW films, however, the significance of the surface and substrate effects are modified, probably due to the effect of chain stiffness on the Tg and different degrees of substrate segregation of the higher s% chains at different MW’s, respectively.

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film can decrease if its thickness is decreased to be comparable to ξ. On the basis of this idea, diminishment of ξ, known to take place at temperatures sufficiently high above the bulk Tg, Tg,bulk, or with addition of an antiplasticizer,35 was used to explain the flattening of the thickness, h, dependences of the Tg observed in polystyrene (PS) supported films or stacked films at high cooling rates19,21,47 and of a variety of supported films upon addition of an antiplasticizer.48,50,56 The idea that variations in the Tg(h) of polymer films might be caused by surface effects was proposed by Keddie et al.27 These workers observed that the Tg(h) (or more precisely, T g (h)/T g (∞)) of PS films supported by silicon was independent of the polymer molecular weight (MW = 120 to 3,000 kg/mol) and therefore could not be caused by physical confinement of the polymer chains, namely the radius of gyration, Rg, of the polymer falling below the film thickness upon thickness reduction. Subsequent studies57,58 extending the measurement to PS films with lower MWs below the entanglement MW (Me ∼ 18 kg/mol for PS59) still found that the Tg(h)/Tg(∞) was independent of the MW. Because of these results, it had been perceived for some time that Tg(h)/ Tg(∞) might not depend on MW. More recently, the Vogt group60 found that the Tg of atatic PMMA (a-PMMA) films supported by silica (a-PMMA/SiOx) with number-average molecular weights, Mn, between 31.7 and 46 kg/mol could change from being enlarged to depressed if the polydispersity index (PDI) of the polymer was increased sufficiently to include oligomers in the films. A subsequent

INTRODUCTION Since Jackson and McKenna1 discovered that the glass transition temperature, Tg, of o-terphenyl and benzyl alcohol, upon confinement in nanoporous glasses, can be visibly different from the bulk Tg, many other glass-formers,2,3 notably polymeric thin films4−27 had been found to exhibit similar and related phenomena. These and other ensuing experimental and theoretical investigations28−39 have now forged a convergent view about the physical underpinning of the observations. Typically, changes in the Tg are interpreted to be caused by competitions27 between effects of the free surface that tend to lower the Tg17 and those of the substrate surface that tend to increase the Tg or cause no change to it.17,40 Consistent with this view, capped films (namely the films that are sandwiched between solid substrates) where the polymer interacts favorably with the substrate surface had been found to exhibit Tg enhancement.41 However, for the capped films where the polymer interacts weakly with the substrate surface, the slow equilibration of the polymer chains when they are in contact with a substrate surface42−44 (attributable to their tendency to adsorb irreversibly onto a solid substrate even when the polymer−substrate interactions are weak42), may cause the Tg of these films to be depressed at first upon fabrication, but slowly approach the bulk value over time.45,46 There are also occasions where “physical confinement” is needed to explain the experimental observations.19,21,35,47−50 It has been shown by experiments51,52 and simulations53,54 that transition to the glassy state occurs when a net change in the molecular arrangements of a specimen necessitates cooperative movements of a cluster of molecules called cooperativity (volume ≈ ξ3),55 which takes an average amount of time equals to the experimental time to complete. Accordingly, the Tg of a © 2016 American Chemical Society

Received: January 16, 2016 Revised: February 16, 2016 Published: March 21, 2016 2671

DOI: 10.1021/acs.macromol.6b00108 Macromolecules 2016, 49, 2671−2678

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Macromolecules

PMMA2.5K; 84.6 ± 0.5 °C for s-PMMA2.5K; 118.1 ± 0.5 °C for aPMMA50K, which is the a-PMMA with Mn = 47 kg/mol; and 125 ± 1 °C for s-PMMA50K, which is the s-PMMA with Mn = 52 kg/mol) for up to 2 h in a ∼ 10−2 Torr vacuum administered by an oil-free pump to relax the films and remove residual solvent. The preannealing temperatures applied to the films are listed in Table S1 (Supporting Information). We find that the preannealing conditions employed are effective in inhibiting dewetting of the films while at the same time providing adequate thermal pretreatment of the films to yield good reproducibility of the Tg measurement in repeated runs. Bilayer films, consisting of PMMAs with similar MW but different tacticities, were used to determine relative sensitivity of different PMMAs to the Tg perturbations by different interfaces. To make the bilayer films, the PMMA used to make the bottom layer was first spun cast onto the substrate to produce a film with thickness h > 4Rg. Then the film was annealed at >20 oC above the bulk Tg in vacuum for 10 h. Afterward, the film was rinsed thoroughly multiple times in fresh toluene. Table 1 shows the average residual film thickness of the

experiment of Lan and Torkelson61 found that for a-PMMA (and also poly(1-ethyl cyclopentyl methacrylate)) supported by silica with narrower MW distributions, the Tg of the films with Mn > 180 kg/mol were enlargedconsistent with the original result of Grohens et al.8but the Tg of the films with Mn < 4.5 kg/mol were depressedconsistent with the aforementioned result of the Vogt group. In a later experiment by us,62 it was found that the Tg of poly(α-methyl styrene) (PαMS) supported by silica with Mn ≥ 20 kg/mol was depressed, but the Tg of the films with Mn = 1.3 kg/mol was independent of the film thickness. These results of different groups show that there is no universal correlation between ΔTg(h) (Tg(h) − Tg(∞)) and the polymer MW when ΔTg(h) exhibits dependence on the MW. What we find surprising is the sensitivity of the noted MW dependence of ΔTg(h) on the polymer tacticity in the case of PMMA supported by silica (PMMA/SiOx), which constitutes the motivation of this work. In an experiment by a former group member,63 the Tg of PMMA/SiOx with Mn = 2.5 kg/mol and PDI = 1.09 was found to be enlarged, which ostensibly contradicts the result of Lan and Torkelson mentioned above. A closer examination reveals that the syndiotacticity (s%) of the PMMAs used in different experiments were different. In the former,63 s% = 69% (iso:hetero:syndio = i:h:s = 7:24:69), but in Lan et al.’s61 it was 59% (i:h:s = 3:38:59). In the following, we detail a comprehensive investigation undertaken to understand the different results. In particular, we measured the Tg of PMMA/SiOx with two different MW’s (Mn = 2.5 and ∼50 kg/mol) and two different s%’s (= 56% and 69 or 79%) similar to those studied in refs 61 and 63. Besides the usual PMMA/SiOx single-layer films, we also studied bilayer films containing PMMAs with similar MW but different s% supported by SiOx. The results enable us to identify which PMMA is more susceptible to Tg perturbation by which interface (namely the free surface or substrate interface) and thereby envisage the physical underpinning of the variability of the Tg of PMMA/SiOx with MW and tacticity.

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Table 1. Average Residual Layer Thickness of the PMMAs Studied films a-PMMA2.5K/SiOx s-PMMA2.5K/SiOx a-PMMA47K/SiOx s-PMMA52K/SiOx

average residual thickness, hres (nm) 0.3 0.5 1.1 3.7

± ± ± ±

0.1 0.1 0.1 0.2

PMMAs studied. Afterward the second, upper polymer layer was spuncast onto the residual film of the first layer to form a bilayer film with total thickness, h. All the bilayer films were annealed under the same preannealing conditions as the corresponding single-layer films listed in Table S1 before Tg measurements. Thin Film Tg Measurement by Ellipsometry. A singlewavelength (λ = 633 nm) Stokes ellipsometer by Gaertner Scientific Corp. was used to measure the Tg of the films. A homemade sample stage equipped with a Peltier element connected to a temperature controller was used to control the sample temperature. In this experiment, the ellipsometric angles, Δ and Ψ, were measured as the sample temperature, T, was ramped down at rates = 2 °C/min. At the glass transition, the thermal expansion coefficient of the film changes steadily from a higher rubbery-state value to a lower glassy-state value. Assuming the ellipsometric angles to be proportional to the film thickness (which is expected to hold since the film thickness change should be only a few percent given the typical thermal expansivity of polymer films to be on the order of ∼10−4 °C−1 6,13), the Tg can be determined by plotting the first derivative of Δ or Ψ versus T, then fitting the resulting curve to separate straight lines in the glassy and liquid states and identifying the midpoint of the transition region to be the Tg. Figure 1 illustrates this measurement scheme. We limited our study to films with thickness ≥ ∼ 20 nm as we found that the Tg measurement became unreliable below this thickness. To calibrate the temperature of our measurement setup, small pieces of Gallium (melting temperature, Tm = 29.7 °C) and indium (Tm = 159.6 °C) were placed on a silicon substrate, which was then put on top of the sample stage. We observed when these specimens began to melt under an optical microscope as the set-point temperature was slowly increased past the melting points. This permitted us to relate the sample temperature with the setpoint temperature at the two melting temperatures. To decide how to interpolate between and extrapolate from these points, we undertook a second calibration procedure by attaching a thermocouple to the top of the sample stage with thermal grease. We found that the calibrations obtained by the two methods agreed within 0.7 °C. Moreover, the sample temperature varied linearly with the set-point temperature throughout the temperature range examined.

MATERIALS AND METHODS

Sample Preparation. The substrates used were single crystal (001) silicon wafers covered by a 102 ± 2 nm thick thermal oxide, purchased from Si-Tech, Inc. Before use, the wafers were cut into 1 cm × 1 cm slides and cleaned in a mixture of sulfuric acid and hydrogen peroxide in a 7:3 volume ratio at 130 °C for 20 min. Then the substrates were thoroughly rinsed in deionized water and dried by 99.99% nitrogen gas. This was followed by further cleaning in oxygen plasma for 20 min, whereupon the slides were ready for use. The PMMA polymers were purchased from Polymer Source, Inc. (Montreal, Canada), with Mn = 2.5 (i:h:s = 6:38:56 and 7:24:69), 47.3 (i:h:s = 6:38:56) and 52 kg/mol (i:h:s = 2:19:79) and PDI less than 1.1. Typically, PMMAs with s% ≤ 59% were referred to as atatic (aPMMA)60,61 and those with s% ≥ 80% syndiotactic (s-PMMA).8,9 In this paper, we relax the condition for the latter8,9 and refer to the PMMAs with s% ≥ 69% s-PMMA. To make a film, we dissolved the as-purchased polymer in toluene with various concentrations between 0.75 and 5 wt %. Then we filtered the solution through a Teflon filter with nominal pore size of 0.1 μm before spin-coating it onto a cleaned substrate whereupon a thin film of the polymer was formed. The thickness of the polymer film was controlled by controlling the concentration of the polymer solution and spinning speed, and measured by ellipsometry at three different locations of the film. Deviation in the film thickness was 79% as syndiotactic and apparently those with s% < 79% and i% < ∼10% as atactic. For ease of comparison, the Tg(Mn) data were normalized by the respective value of Tg(Mn = ∞) before being plotted versus 1/Mn in Figure 4. The data show that the

Figure 3. Plots of Tg(h) − Tg(∞)) versus total film thickness, h, of the Mn ≈ 50 kg/mol a-PMMA and s-PMMA single-layer (solid symbols) and bilayer films (open symbols). The solid curves are guides to the eye. Schematic representations of different films are shown by the drawings, where a-PMMA50K is labeled by “a” (white) and sPMMA50K by “s” (black), and the substrate is marked by diagonal stripes.

previous published results.8,9,27,64 With these findings, the signs of the ΔTg(h) are reversed with respect to those of the 2.5 kg/ mol films in Figure 2. The possible reasons for this sign reversal will be discussed in the next section. Now we focus on the remaining data of Figure 3, namely the data obtained from the bilayer films. For the a-PMMA/s-PMMA/SiOx films (open squares), insertion of the s-PMMA layer beneath the a-PMMA layer promotes the Tg enhancement of the a-PMMA layer. Similarly, for the s-PMMA/a-PMMA/SiOx films (open triangles), insertion of the a-PMMA layer beneath the sPMMA layer promotes the Tg depression of the s-PMMA layer. These results indicate that the presence of s-PMMA next to the silica substrate enhances the Tg of the films, in keeping with the results found in Figure 2. Given this, to fully account for the result of Figure 3 would require that the perturbations from the free surface, which are presumed to bring about depression of the surface Tg, are stronger in s-PMMA50K than in aPMMA50K. This is ostensibly opposite to what we concluded above about the effect of the free surface on the 2.5 kg/mol films. In the next section, we shall address this and the other apparently contradictory observations, and more generally the possible reasons for the MW dependence revealed by Figures 2 and 3 of the effects of the interfaces on the ΔTg of the a- and sPMMA films.

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Figure 4. Plots of Tg(Mn)/Tg(Mn=∞) versus 105/Mn using the data of a-PMMA and s-PMMA (with s% > 79) provided by our polymer supplier and the data of Beevers et al.68 The normalization constant, Tg(Mn=∞) is 386 K (i.e., 113 °C) for a-PMMA (supplier), 405 K (i.e., 132 °C) for s-PMMA (with s% > 79%, supplier), and 390 K (i.e., 117 °C) for the data of Beevers et al. The solid line is a fit to Fox and Flory’s relation69 and the dashed line is the published fit to Gibbs’ equation67 by Beevers et al.

DISCUSSION First, we consider the reason why the surface Tg of s-PMMA may be less depressed at Mn = 2.5 kg/mol than at Mn ∼ 50 kg/ mol. In a recent experiment,62 we observed a similar effect in poly(α-methyl styrene) supported by silica (PαMS/SiOx). There, we attributed the observation to flattening of the MW dependence of the Tg of the PαMS bulk polymer at low MW as noted by Cowie et al.,65 which could imply that the Tg of PαMS becomes less sensitive to changes in free volume at low MW. In other words, if depression of the surface Tg is caused by augmentation of free volume at the free surface, the surface Tg should be little modified from the bulk value at low MW. Cowie et al.65 also found that the flattening of the Tg(MW) dependence of bulk PαMS could be accounted for by the model of Gibbs and DiMarzio66 (or more precisely the original equation derived by Gibbs67). In the Gibbs and DiMarzio model, the glass transition is conjectured to be caused by the

data of s-PMMA overlaps with that of Beevers et al., which fits well to Gibbs’ equation (dashed line, Figure 4) and saturates toward a constant value of Tg(Mn)/Tg(Mn=∞) ≈ 0.86 at low Mn; on the other hand, the data of a-PMMA fits better to Fox and Flory’s equation69 and shows no clear tendency to saturate. There are, however, caveats with the data of Figure 4. Most notably, they had been collected from PMMAs with a range of tacticities. Therefore, variations of the Tg with tacticity can contribute to the apparent Mn dependence seen in Figure 4. In 2674

DOI: 10.1021/acs.macromol.6b00108 Macromolecules 2016, 49, 2671−2678

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Macromolecules addition, although the Tg(Mn)/Tg(Mn=∞) data of Beevers et al. (open circles, Figure 4) agrees better with the supplier’s data for s-PMMA (open triangles, Figure 4), their Tg(Mn=∞) value actually agrees better with that of the supplier’s value for aPMMA than for s-PMMA. With all the caveats noted, we however also perceive credibility of the data of Figure 4. First, different data sets show the expected relative trends for the Tg(Mn) dependence given that s-PMMA has a larger value for the conformational energy, ε, than a-PMMA does. In addition, the saturation value of Tg(Mn) ≈ 0.86 Tg(Mn=∞) at low Mn for the data of Beevers et al. (dashed line, Figure 4) is approximately equal to 1.3 ε/kB based on the value of ε = 879 cal/mol measured by Grohens et al.9 for bulk s-PMMA. In contrast, the value of ε/kB for bulk a-PMMA is much lower, equal to 237 K (or −36 °C),9 which would correspond to a saturation value of Tg(Mn)/Tg(Mn=∞) ≈ 0.7 in Figure 4. Such a low saturation value corroborates with the data of a-PMMA in Figure 4 showing no obvious tendency for saturation. Lastly, by using the bulk Tg data reported above, one finds the value of Tg(Mn)/Tg(Mn=∞) to be 0.92 and 0.88 for, respectively, the aPMMA2.5K and s-PMMA2.5K polymers studied here. These values fall well into the respective curves according to the nominal tacticities of these polymers. With the credibility of the data of Figure 4 noted, we proceed to use this figure to account for the ΔTg data of Figures 2 and 3. Figure 4 shows that at Mn = 2.5 kg/mol, corresponding to 105/Mn = 40 (kg/mol)−1, the Tg of s-PMMA should approach saturation, but that of aPMMA not. Therefore, perturbations from the free-surface, if operative by free-volume augmentation, can cause depression of the surface Tg in a-PMMA2.5K but not in s-PMMA2.5K. This, together with the stronger interactions s-PMMA has with the substrate than a-PMMA does, the data of Figure 2 can be explained. Now we address the sign reversal of the ΔTg’s in going between different MWs. First we address the sign reversal in the a-PMMA films. From Figure 4, the Tg(Mn) dependence of aPMMA follows the Fox and Flory equation and does not show any tendency to saturate at low Mn, which is similar to the Tg(Mn) dependence found of bulk polystyrene (PS).57 As mentioned above, PS films show no MW dependence in the Tg(h)/Tg(∞) measurements,57 implying that perturbation from the free surface to Tg(h)/Tg(∞) is independent of MW. Since a-PMMA and PS exhibit similar Tg(Mn) dependences, one may expect the same for a-PMMA. Henceforth, the sign change, from negative to positive, of the ΔTg of the a-PMMA films in going from low to high MW must be caused by an increase in the polymer−substrate interaction with MW. We tentatively suggest this to be caused by a tendency of the polymer chains with statistically higher s% to segregate to the substrate surface. We expect this tendency to be higher with higher MW because firstly enthalpy gain per chain is higher at higher MW’s and secondly entropy factors, which tend to mix different components together, are more significant at lower MW’s. The idea that the tacticity of the PMMA chains being adsorbed to the substrate may influence the ΔTg of the films is readily evident from the visibly different ΔTg(h)’s found above between the a-PMMA2.5K single- and bilayer films. On the other hand, we notice that the ΔTg(h) of the a-PMMA50K single-layer films is bigger than that of the a-PMMA2.5K/sPMMA2.5K/SiOx bilayer films. One possible reason for this observation is that the amount of perturbations the adsorbed layer can exercise on a film’s Tg depends on its thickness,

especially how it compares to the size of cooperativity at the glass transition. A study is underway to address this question. Finally, we address the sign change of the ΔTg of the sPMMA films, which goes from positive to negative as the MW is increased from 2.5 to ∼50 kg/mol. We believe that the main driver for this phenomenon is the saturation of the Tg of sPMMA with free volume change near Mn = 2.5 kg/mol (Figure 4), which limits the surface Tg of the s-PMMA2.5K from getting depressed, but not for the s-PMMA50K polymer. The facts that both the Tg’s of the a-PMMA50K single- and bilayer films are enhanced while those of the s-PMMA50K films are depressed, and s-PMMA interacts more strongly with the substrate surface than a-PMMA does imply that perturbation from the free surface on the Tg is stronger in the s-PMMA50K than in the a-PMMA50K polymer. Dhinojwala et al.70 found that the orientation order of the phenyl groups in PS supported by sapphire was higher at the interfaces. A higher structural order has been attributed to stronger glassy behaviors, namely slower rise in the alpha relaxation time or viscosity with cooling at the Tg.71 We ponder if the bigger depression in the surface Tg implied in our result of the s-PMMA50K polymer relative to that of the a-PMMA50K polymer may be caused by the more regular microstructure of s-PMMA as that could enable better orientation order at the free surface. According to Ngai et al.’s coupling model description of glassy dynamics,72−75 stronger glassy behavior corresponds to a smaller coupling parameter, n, which physically means weaker coupling between different simpler, one-molecule (primitive) relaxations, and better molecular orientation of polymer chains had been noted to corroborate with faster segmental dynamics.72 Presumably any net orientation order at the surface tapers toward zero within roughly the same distance of the order of a monomer thickness.70 Then reduction of the surface Tg should be stronger in s-PMMA than in a-PMMA, in keeping with the above observation. Some time ago, Ngai had proposed an explanation for the different ΔTg’s found of the high-Mw sPMMA and i-PMMA films by using the coupling model.73 It is not immediately obvious how it can be adapted to account for the MW and tacticity dependences found here of the ΔTg. It is hoped that presentation of the proposed explanation herein may stimulate further discussions. Lastly, we compare the (chain stiffness) effect proposed here, relating the flattening of the Tg(MW) dependence at low MW (Figure 4) with suppression of reduction in the surface Tg, with the effect proposed by Zhang et al.76 in explaining the lack of Tg confinement effect found of cyclic PS (c-PS) supported by silicon. While the bulk Tg of c-PS was also noted to be independent of MW for Mn = 3.4 and 9.1 kg/mol, Zhang et al. attributed the flat Tg(h) dependence to a speculated invariance of the molecular packing frustration at the free surface relative to that in the bulk polymer. Such invariance was inferred from the bulk fragility of a low-MW c-PS (Mn = 4.6 kg/mol) being similar to the bulk fragility of linear PS in the Mn → ∞ limit, and the assumption that fragility is a reflection of packing frustration.77 Since the (chain-stiffness) effect considered here concerns whether free volume is able to cause changes to the Tg, it is clearly different from that considered by Zhang et al.

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CONCLUSION A systematic study was performed on the thickness, h, dependence of the Tg of PMMA films supported by silica with two different MW (Mn = 2.5 and ∼50 kg/mol) and syndiotacticities (s% = 56% for a-PMMA and 69 or 79% for s2675

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Macromolecules PMMA). Our result shows that the sign of ΔTg(h) is sensitive to both the MW and tacticity. To determine which interface (air or substrate) dominates in which of the four kinds of PMMA films studied, we constructed bilayer films consisting of PMMAs with similar MW but different s%. We find that while the sign of ΔTg(h) is consistent with the convention that it is governed by an interplay between two competing surface effects−namely, the free surface, which depresses the Tg and the substrate surface, which enhances it−the MW and tacticity modulate these effects in nontrivial ways. Specifically, our result implies that (1) the substrate effect on the Tg is stronger in sPMMA than in a-PMMA and (2) at Mn = 50 kg/mol, the surface Tg of s-PMMA is depressed more from the bulk Tg than is that of a-PMMA, but at Mn = 2.5 kg/mol the effect of the free surface becomes ineffective in s-PMMA. We contemplate this ineffectiveness of the free surface to be brought about by the emergent effect of chain stiffness on the Tg of polymers as discussed by Gibbs and DiMarzio. In addition, preferential segregation of the statistically higher s% polymer chains to the substrate surface in the a-PMMA films may also play a role.

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(8) Grohens, Y.; Brogly, M.; Labbe, C.; David, M. O.; Schultz, J. Glass Transition of Stereoregular Poly(Methyl Methacrylate) at Interfaces. Langmuir 1998, 14 (11), 2929−2932. (9) Grohens, Y.; Hamon, L.; Reiter, G.; Soldera, A.; Holl, Y. Some Relevant Parameters Affecting the Glass Transition of Supported Ultra-Thin Polymer Films. Eur. Phys. J. E: Soft Matter Biol. Phys. 2002, 8, 217−224. (10) Kim, J.-H.; Jang, J.; Zin, W.-C. Estimation of the Thickness Dependence of the Glass Transition Temperature in Various Thin Polymer Films. Langmuir 2000, 16, 4064−4067. (11) Fukao, K.; Uno, S.; Miyamoto, Y.; Hoshino, A.; Miyaji, H. Dynamics of Alpha and Beta Processes in Thin Polymer Films: Poly(Vinyl Acetate and Poly(Methyl Methacrylate). Phys. Rev. E: Stat. Phys., Plasmas, Fluids, Relat. Interdiscip. Top. 2001, 64, 051807. (12) Fukao, K. Dynamics in Thin Polymer Films by Dielectric Spectroscopy. Eur. Phys. J. E: Soft Matter Biol. Phys. 2003, 12, 119− 125. (13) Tsui, O. K. C.; Russell, T. P.; Hawker, C. J. Effect of Interfacial Interactions on the Glass Transition of Polymer Thin Films. Macromolecules 2001, 34, 5535−5539. (14) Xie, F.; Zhang, H. F.; Lee, F. K.; Du, B.; Tsui, O. K. C.; Yokoe, Y.; Tanaka, K.; Takahara, A.; Kajiyama, T.; He, T. Effect of Low Surface Energy Chain Ends on the Glass Transition Temperature of Polymer Thin Films. Macromolecules 2002, 35, 1491−1492. (15) Serghei, A.; Tress, M.; Kremer, F. Confinement Effects on the Relaxation Time Distribution of the Dynamic Glass Transition in Ultrathin Polymer Films. Macromolecules 2006, 39, 9385−9387. (16) Prevosto, D.; Napolitano, S.; Pingue, P.; Capaccioli, S.; Lucchesi, M. Investigation of Structural Relaxation and Surface Modification of Ultrathin Films of Poly(Ethylene Terephthalate). Eur. Phys. J.: Spec. Top. 2007, 141, 193−198. (17) Ellison, C. J.; Torkelson, M. The Distribution of GlassTransition Temperatures in Nanoscopically Confined Glass Formers. Nat. Mater. 2003, 2, 695−670. (18) Priestley, R. D.; Mundra, M. K.; Barnett, N. J.; Broadbelt, L. J.; Torkelson, J. M. Effects of Nanoscale Confinement and Interfaces on the Glass Transition Temperatures of a Series of Poly(NMethacrylate) Films. Aust. J. Chem. 2007, 60, 765−771. (19) Fakhraai, Z.; Forrest, J. A. Probing Slow Dynamics in Supported Thin Polymer Films. Phys. Rev. Lett. 2005, 95, 025701. (20) Fakhraai, Z.; Forrest, J. A. Measuring the Surface Dynamics of Glassy Polymers. Science 2008, 319, 600−604. (21) Koh, Y. P.; Simon, S. L. Structural Relaxation of Stacked Ultrathin Polystyrene Films. J. Polym. Sci., Part B: Polym. Phys. 2008, 46, 2741−2753. (22) Paeng, K.; Swallen, S. F.; Ediger, M. D. Direct Measurement of Molecular Motion in Freestanding Polystyrene Thin Films. J. Am. Chem. Soc. 2011, 133, 8444−8447. (23) Zhu, L.; Brian, C. W.; Swallen, S. F.; Straus, P. T.; Ediger, M. D.; Yu, L. Surface Self-Diffusion of an Organic Glass. Phys. Rev. Lett. 2011, 106 (25), 256103. (24) Pye, J. E.; Roth, C. B. Two Simultaneous Mechanisms Causing Glass Transition Temperature Reductions in High Molecularweight Freestanding Polymer Films as Measured by Transmission Ellipsometry. Phys. Rev. Lett. 2011, 107, 235701. (25) Glynos, E.; Frieberg, B.; Oh, H.; Liu, M.; Gidley, D. W.; Green, P. F. Role of Molecular Architecture on the Vitrification of Polymer Thin Films. Phys. Rev. Lett. 2011, 106, 128301. (26) Yin, H.; Napolitano, S.; Schonhals, A. Molecular Mobility and Glass Transition of Thin Films of Poly(Bisphenol a Carbonate). Macromolecules 2012, 45, 1652−1662. (27) Keddie, J. L.; Jones, R. A. L.; Cory, R. A. Interface and Surface Effects on the Glass-Transition Temperature in Thin Polymer Films. Faraday Discuss. 1994, 98, 219−230. (28) de Gennes, P.-G. Glass Transitions in Thin Polymer Films. Eur. Phys. J. E: Soft Matter Biol. Phys. 2000, 2, 201−205. (29) Long, D.; Lequeux, F. Heterogeneous Dynamics at the Glass Transition in Van Der Waals Liquids, in Bulk and in Thin Films. Eur. Phys. J. E: Soft Matter Biol. Phys. 2001, 4, 371−387.

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.6b00108. Table S1, containing the preannealing conditions used in this study (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected] (O.K.C.T.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We are grateful for the support of the National Science Foundation (DMR-1310536).

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REFERENCES

(1) Jackson, C. L.; McKenna, G. B. The Glass Transition of Organic Liquids Confined to Pores. J. Non-Cryst. Solids 1991, 131−133, 221− 224. (2) Arndt, M.; Stannarius, R.; Groothues, H.; Hempel, E.; Kremer, F. Length Scale of Cooperativity in the Dynamic Glass Transition. Phys. Rev. Lett. 1997, 79, 2077−2080. (3) Hunter, G. L.; Weeks, E. R. The Physics of the Colloidal Glass Transition. Rep. Prog. Phys. 2012, 75, 066501. (4) Forrest, J. A.; Dalnoki-Veress, K.; Stevens, J. R.; Dutcher, J. R. Effect of Free Surfaces on the Glass Transition Temperature of Thin Polymer Films. Phys. Rev. Lett. 1996, 77, 2002−2005. (5) Sharp, J. S.; Forrest, J. A. Dielectric and Ellipsometric Studies of the Dynamics in Thin Films of Isotactic Poly(Methylmethacrylate) with One Free Surface. Phys. Rev. E: Stat. Phys., Plasmas, Fluids, Relat. Interdiscip. Top. 2003, 67, 031805. (6) van Zanten, J. H.; Wallace, W. E.; Wu, W.-L. Effect of Strongly Favorable Substrate Interactions on the Thermal Properties of Ultrathin Polymer Films. Phys. Rev. E: Stat. Phys., Plasmas, Fluids, Relat. Interdiscip. Top. 1996, 53, R2053−R2056. (7) Tanaka, K.; Takahara, A.; Kajiyama, T. Effect of Polydispersity on Surface Molecular Motion of Polystyrene Films. Macromolecules 1997, 30, 6626−6632. 2676

DOI: 10.1021/acs.macromol.6b00108 Macromolecules 2016, 49, 2671−2678

Article

Macromolecules (30) Lipson, J. E. G.; Milner, S. T. Local and Average Glass Transitions in Polymer Thin Films. Macromolecules 2010, 43, 9874− 9880. (31) Milner, S. T.; Lipson, J. E. G. Delayed Glassification Model for Free-Surface Suppression of Tg in Polymer Glasses. Macromolecules 2010, 43, 9865−9873. (32) Tito, N. B.; Lipson, J. E. G.; Milner, S. T. Lattice Model of Mobility at Interfaces: Free Surfaces, Substrates, and Bilayers. Soft Matter 2013, 9, 9403−9413. (33) Varnik, F.; Baschnagel, J.; Binder, K. Reduction of the Glass Transition Temperature in Polymer Films: A Molecular-Dynamics Study. Phys. Rev. E: Stat. Phys., Plasmas, Fluids, Relat. Interdiscip. Top. 2002, 65, 021507. (34) Hanakata, P. Z.; Douglas, J. F.; Starr, F. W. Interfacial Mobility Scale Determines the Scale of Collective Motion and Relaxation Rate in Polymer Films. Nat. Commun. 2014, 5, 4163. (35) Riggleman, R. A.; Yoshimoto, K.; Douglas, J. F.; de Pablo, J. J. Influence of Confinement on the Fragility of Antiplasticized and Pure Polymer Films. Phys. Rev. Lett. 2006, 97, 045502. (36) Shavit, A.; Riggleman, R. A. Influence of Backbone Rigidity on Nanoscale Confinement Effects in Model Glass-Forming Polymers. Macromolecules 2013, 46, 5044−5052. (37) Lam, C.-H.; Tsui, O. K. C. Crossover to Surface Flow in Supercooled Unentangled Polymer Films. Phys. Rev. E 2013, 88 (88), 042604. (38) Lang, R. J.; Simmons, D. S. Interfacial Dynamic Length Scales in the Glass Transition of a Model Freestanding Polymer Film and Their Connection to Cooperative Motion. Macromolecules 2013, 46, 9818− 9825. (39) Mirigian, S.; Schweizer, K. S. Slow Relaxation, Spatial Mobility Gradients, and Vitrification in Confined Films. J. Chem. Phys. 2014, 141, 161103. (40) Roth, C. B.; McNerny, K. L.; Jager, W. F.; Torkelson, J. M. Eliminating the Enhanced Mobility at the Free Surface of Polystyrene: Fluorescence Studies of the Glass Transition Temperature in Thin Bilayer Films of Immiscible Polymers. Macromolecules 2007, 40, 2568−2574. (41) Rittigstein, P.; Priestley, R. D.; Broadbelt, L. J.; Torkelson, J. M. Model Polymer Nanocomposites Provide an Understanding of Confinement Effects in Real Nanocomposites. Nat. Mater. 2007, 6, 278−282. (42) Fujii, Y.; Yang, Z.; Leach, J.; Atarashi, H.; Tanaka, K.; Tsui, O. K. C. Affinity of Polystyrene Films to Hydrogen-Passicated Silicon and Its Relevance to the Tg of the Films. Macromolecules 2009, 42, 7418− 7422. (43) Housmans, C.; Sferrazza, M.; Napolitano, S. Kinetics of Irreversible Chain Adsorption. Macromolecules 2014, 47, 3390−3393. (44) Jiang, N.; Shang, J.; Di, X.; Endoh, M. K.; Koga, T. Formation Mechanism of High-Density, Flattened Polymer Nanolayers Adsorbed on Planar Solids. Macromolecules 2014, 47, 2682−2689. (45) Napolitano, S.; Wubbenhorst, M. The Lifetime of the Deviations from Bulk Behaviour in Polymers Confined at the Nanoscale. Nat. Commun. 2011, 2, 260. (46) Napolitano, S.; Capponi, S.; Vanroy, B. Glassy Dynamics of Soft Matter under 1d Confinement: How Irreversible Adsorption Affects Molecular Packing, Mobility Gradients and Orientational Polarization in Thin Films. Eur. Phys. J. E: Soft Matter Biol. Phys. 2013, 36, 61−97. (47) Efremov, M. Y.; Olson, E. A.; Zhang, M.; Zhang, Z.; Allen, L. H. Glass Transition in Ultrathin Polymer Films: Calorimetric Study. Phys. Rev. Lett. 2003, 91, 085703. (48) Ellison, C. J.; Ruszkowski, R. L.; Fredin, N. J.; Torkelson, J. M. Dramatic Reduction of the Effect of Nanoconfinement on the Glass Transition of Polymer Films Via Addition of Small-Molecule Diluent. Phys. Rev. Lett. 2004, 92, 095702. (49) Mundra, M. K.; Ellison, C. J.; Rittigstein, P.; Torkelson, J. M. Fluorescence Studies of Confinement in Polymer Films and Nanocomposites: Glass Transition Temperature, Plasticizer Effects, and Sensitivity to Stress Relaxation and Local Polarity. Eur. Phys. J.: Spec. Top. 2007, 141, 143−151.

(50) Kim, S.; Mundra, M. K.; Roth, C. B.; Torkelson, J. M. Suppression of the Tg-Nanoconfinement Effect in Thin Poly(Vinyl Acetate) Films by Sorbed Water. Macromolecules 2010, 43, 5158− 5161. (51) Tracht, U.; Wilhelm, M.; Heuer, A.; Feng, H.; Schmidt-Rohr, K.; Spiess, H. W. Length Scale of Dynamic Heterogeneities at the Glass Transition Determined by Multidimensional Nuclear Magnetic Resonance. Phys. Rev. Lett. 1998, 81, 2727−2730. (52) Vidal Russell, E. V.; Israeloff, N. E. Direct Observation of Molecular Cooperativity near the Glass Transition. Nature 2000, 408, 695−8. (53) Donati, C.; Douglas, J. F.; Kob, W.; Plimpton, S. J.; Poole, P. H.; Glotzer, S. C. Stringlike Cooperative Motion in a Supercooled Liquid. Phys. Rev. Lett. 1998, 80, 2338−2341. (54) Hanakata, P. Z.; Douglas, J. F.; Starr, F. W. Local Variation of Fragility and Glass Transition Temperature of Ultra-Thin Supported Polymer Films. J. Chem. Phys. 2012, 137, 244901. (55) Adam, G.; Gibbs, J. H. The Temperature Dependence of Cooperative Relaxation in Glass-Forming Liquids. J. Chem. Phys. 1965, 43, 139−146. (56) Nguyen, H. K.; Labardi, M.; Lucchesi, M.; Rolla, P.; Prevosto, D. Plasticization in Ultrathin Polymer Films: The Role of Supporting Substrate and Annealing. Macromolecules 2013, 46, 555−561. (57) Tsui, O. K. C.; Zhang, H. F. Effect of Chain Ends and Chain Entanglement on Glass Transition Temperature of Polymer Thin Films. Macromolecules 2001, 34, 9139−9142. (58) Ellison, C. J.; Mundra, M. K.; Torkelson, J. M. Impacts of Polystyrene Molecular Weight and Modification to the Repeat Unit Structure on the Glass Transition-Nanoconfinement Effect and the Cooperativity Length Scale. Macromolecules 2005, 38, 1767−1778. (59) Fetters, L. J.; Lohse, D. J.; Milner, S. T.; Graessley, W. W. Packing Length Influence in Linear Polymer Melts on the Entanglement, Critical, and Reptation Molecular Weights. Macromolecules 1999, 32, 6847−6851. (60) Dion, M.; Larson, A. B.; Vogt, B. D. Impact of Low-Molecular Mass Components (Oligomers) on the Glass Transition in Thin Films of Poly(Methyl Methacrylate). J. Polym. Sci., Part B: Polym. Phys. 2010, 48, 2366−2370. (61) Lan, T.; Torkelson, J. M. Methacrylate-Based Polymer Films Useful in Lithographic Applications Exhibit Different Glass Transition Temperature-Confinement Effects at High and Low Molecular Weight. Polymer 2014, 55, 1249−1258. (62) Geng, K.; Chen, F.; Tsui, O. K. C. Molecular-Weight Dependent Tg Depression of Silica-Supported Poly(Alpha-Methyl Styrene) Films. J. Non-Cryst. Solids 2015, 407, 296−301. (63) Li, R. N.; Chen, F.; Lam, C.-H.; Tsui, O. K. C. Viscosity of Polymethylmethacrylate on Silica: Epitome of Systems with Strong Polymer−Substrate Interactions. Macromolecules 2013, 46, 7889− 7893. (64) Tan, A. W.; Torkelson, J. M. Poly(Methyl Methacrylate) Nanotubes in Aao Templates: Designing Nanotube Thickness and Characterizing the Tg-Confinement Effect by Dsc. Polymer 2016, 82, 327−336. (65) Cowie, J. M. G.; Toporowski, P. M. The Dependence of Glass Temperature on Molecular Weight for Poly Alpha-Methyl Styrene. Eur. Polym. J. 1968, 4, 621−625. (66) Gibbs, J. H.; DiMarzio, E. A. Nature of the Glass Transition and the Glassy State. J. Chem. Phys. 1958, 28, 373−383. (67) Gibbs, J. H. Nature of Glass Transition in Polymers. J. Chem. Phys. 1956, 25, 185−6. (68) Beevers, R. B.; White, E. F. T. Physical Properties of Vinyl Polymers. Part 1.Dependence of the Glass-Transition Temperature of Polymethylmethacrylate on Molecular Weight. Trans. Faraday Soc. 1960, 56, 744−752. (69) Fox, T. G.; Flory, P. J. Second-Order Transition Temperatures and Related Properties of Polystyrene. I. Influence of Molecular Weight. J. Appl. Phys. 1950, 21, 581−591. (70) Gautam, K. S.; Schwab, A. D.; Dhinojwala, A.; Zhang, D.; Dougal, S. M.; Yeganeh, M. S. Molecular Structure of Polystyrene at 2677

DOI: 10.1021/acs.macromol.6b00108 Macromolecules 2016, 49, 2671−2678

Article

Macromolecules Air/Polymer and Solid/Polymer Interfaces. Phys. Rev. Lett. 2000, 85 (18), 3854−3857. (71) Angell, C. A. Formation of Glasses from Liquids and Biopolymers. Science 1995, 267, 1924−1935. (72) Ngai, K. L.; Rizos, A. K.; Plazek, D. J. Reduction of the Glass Temperature of Thin Freely Standing Polymer Films Caused by the Decrease of the Coupling Parameter in the Coupling Model. J. NonCryst. Solids 1998, 235−237, 435−443. (73) Ngai, K. L. Interpreting the Dynamics of Nano-Confined GlassFormers and Thin Polymer Films: Importance of Starting from a Viable Theory for the Bulk. J. Polym. Sci., Part B: Polym. Phys. 2006, 44, 2980−2995. (74) Ngai, K. L.; Prevosto, D.; Grassia, L. Viscoelasticity of Nanobubble-Inflated Ultrathin Polymer Films: Justification by the Coupling Model. J. Polym. Sci., Part B: Polym. Phys. 2013, 51, 214− 224. (75) Prevosto, D.; Capaccioli, S.; Ngai, K. L. Origins of the Two Simultaneous Mechanisms Causing Glass Transition Temperature Reductions in High Molecular Weight Freestanding Polymer Films. J. Chem. Phys. 2014, 140, 074903. (76) Zhang, L.; Elupula, R.; Grayson, S. M.; Torkelson, J. M. Major Impact of Cyclic Chain Topology on the Tg-Confinement Effect of Supported Thin Films of Polystyrene. Macromolecules 2016, 49, 257− 268. (77) Santangelo, P. G.; Roland, C. M.; Chang, T.; Cho, D.; Roovers, J. Dynamics near the Glass Temperature of Low Molecular Weight Cyclic Polystyrene. Macromolecules 2001, 34, 9002−9005.

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