Controlling Interfacial Dynamics: Covalent Bonding versus Physical Adsorption in Polymer Nanocomposites Adam P. Holt,*,† Vera Bocharova,*,‡ Shiwang Cheng,‡ Alexander M. Kisliuk,‡ B. Tyler White,‡ Tomonori Saito,‡ David Uhrig,§ J. P. Mahalik,§,∥ Rajeev Kumar,§,∥ Adam E. Imel,⊥ Thusitha Etampawala,⊥ Halie Martin,⊥ Nicole Sikes,# Bobby G. Sumpter,§,∥ Mark D. Dadmun,‡,⊥ and Alexei P. Sokolov‡,⊥ †
Department of Physics and Astronomy and ⊥Department of Chemistry, University of Tennessee, Knoxville, Tennessee 37996, United States ‡ Chemical Sciences Division, §Center for Nanophase Materials Sciences, and ∥Computer Science and Mathematics Division, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831, United States # Department of Chemistry, Columbus State University, Columbus, Georgia 33232, United States S Supporting Information *
ABSTRACT: It is generally believed that the strength of the polymer−nanoparticle interaction controls the modiﬁcation of near-interface segmental mobility in polymer nanocomposites (PNCs). However, little is known about the eﬀect of covalent bonding on the segmental dynamics and glass transition of matrix-free polymer-grafted nanoparticles (PGNs), especially when compared to PNCs. In this article, we directly compare the static and dynamic properties of poly(2-vinylpyridine)/silica-based nanocomposites with polymer chains either physically adsorbed (PNCs) or covalently bonded (PGNs) to identical silica nanoparticles (RNP = 12.5 nm) for three diﬀerent molecular weight (MW) systems. Interestingly, when the MW of the matrix is as low as 6 kg/mol (RNP/Rg = 5.4) or as high as 140 kg/mol (RNP/Rg= 1.13), both small-angle X-ray scattering and broadband dielectric spectroscopy show similar static and dynamic properties for PNCs and PGNs. However, for the intermediate MW of 18 kg/mol (RNP/Rg = 3.16), the diﬀerence between physical adsorption and covalent bonding can be clearly identiﬁed in the static and dynamic properties of the interfacial layer. We ascribe the diﬀerences in the interfacial properties of PNCs and PGNs to changes in chain stretching, as quantiﬁed by self-consistent ﬁeld theory calculations. These results demonstrate that the dynamic suppression at the interface is aﬀected by the chain stretching; that is, it depends on the anisotropy of the segmental conformations, more so than the strength of the interaction, which suggests that the interfacial dynamics can be eﬀectively tuned by the degree of stretchinga parameter accessible from the MW or grafting density. KEYWORDS: polymer nanocomposites, polymer-grafted nanoparticles, glass transition, segmental dynamics, interfacial dynamics, self-consistent ﬁeld theory glass transition temperature, Tg).16−22 However, the diﬀerent magnitude of changes observed in various PNCs has raised questions about the universality of the underlying mechanism of the adsorption process as the crucial process and may suggest it to be system-speciﬁc rather than a multiscale characteristic.23
ecently, it has been suggested that the slowing down of segmental dynamics in heterogeneous or nanostructured materials such as polymer nanocomposites,1−7 block copolymers,8,9 polymer thin ﬁlms,10−12 ionomers,13 and semicrystalline polymers14 share a common origin associated with the dynamic interfaces in these systems.15 While many experimental and theoretical eﬀorts are underway to validate this generalized statement, in polymer nanocomposites (PNCs), the presence of dynamically aﬀected segments, formed via physical adsorption of polymer chains on the nanoparticle’s surface, has been detected experimentally and considered the primary cause of suppression of segmental dynamics (or the © 2016 American Chemical Society
Received: April 13, 2016 Accepted: June 23, 2016 Published: June 23, 2016 6843
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parameter via the MW for PNCs or the grafting density for PGNs.
Nevertheless, based on the current experimental data and theoretical predictions, it is possible to specify the criteria that may be responsible for the dynamic changes at the interfaces in nanostructured materials: (1) the fragility of the polymer,24,25 (2) the interaction strength,2,4,12 or (3) the softness of conﬁnement.15 Some of these criteria are well-studied experimentally, while others require additional experimental validation. For instance, large changes in the segmental dynamics are observed in PNCs with high Tg polymers (above room temperature) that are also relatively fragile glass formers such as poly(vinyl acetate)/silica,19,26 poly(methyl methacrylate)/silica,1,5 or poly(2-vinylpyridine)/silica nanocomposites.1,19−22 The eﬀect of interaction strength on segmental dynamics, however, is seldom studied. The eﬀect of the strongest interaction, such as covalent bonding, on the dynamic interface of matrix-free polymer-grafted nanoparticles (PGNs, where all polymer chains are covalently bound to the nanoparticles) is not well-understood with only a few experimental studies to date.27−29 Even in case of the most commonly studied polymer−nanoparticle systems which interact via hydrogen bonding (this is typically needed to facilitate good NP dispersion), there are many aspects concerning the formation of the interface that are not yet fully understood, where the eﬀect of molecular weight was only recently analyzed in detail.19 Therefore, to advance our understanding of the physical origin of dynamically altered interfaces, the question of how covalent bonding (or tethering) aﬀects the dynamics of the matrix-free grafted polymer chains relative to the dynamics of physically adsorbed chains becomes fundamentally important to rationally design the physical properties of PNCsand other nanostructured polymer materials. In this article, we directly compare the eﬀect of covalent bonding and physical adsorption (induced by strong hydrogen bonds) on the interfacial properties of well-dispersed poly(2vinylpyridine)/silica nanocomposites (PNCs) and matrix-free poly(2-vinylpyridine) grafted from silica systems using broadband dielectric spectroscopy (BDS), temperature-modulated diﬀerential scanning calorimetry (TMDSC), and small-angle Xray scattering (SAXS) techniques corroborated with detailed self-consistent ﬁeld theory (SCFT) calculations. We investigate three molecular weight (MW) systems for both PGNs (chain grafting density = 0.3 chains/nm2) and PNCs with identical silica (RNP = 12.5 nm) concentrations. Our analysis revealed only a weak change in the dynamics of PGNs and PNCs with high MW (140 kg/mol, RNP/Rg = 1.13), most likely due to the relatively low content (∼6 vol % silica) of nanoparticles. However, there is a clear diﬀerence in the properties of the interfacial layer at intermediate MW (18 kg/mol, RNP/Rg = 3.16), where the interfacial dynamics are more strongly suppressed by the covalent bonding than by physical adsorption. Interestingly, we ﬁnd that PGNs and PNCs with low MW (6 kg/mol, RNP/Rg = 5.4) exhibit an identical suppression of dynamics, a large gradient of segmental mobility, and an interfacial region that occupies essentially the entire material despite the magnitude of the interactions. We propose that these eﬀects are caused by anisotropic chain stretching of segments in the interfacial layer, a phenomenon previously suggested by Oyerokun and Schweizer,30 that systematically decreases with increasing MW for PNCs but increases for the PGN systems. Our results suggest that the interfacial dynamics in PNCs or PGNs can be tuned by controlling the degree of stretching at the interface, which should be an adjustable
RESULTS AND DISCUSSION The excellent quality of nanoparticle dispersion in our samples is conﬁrmed by SAXS and transmission electron microscopy images (Figure 1). The changes in macroscopic properties (e.g.,
Figure 1. (a) Small-angle X-ray scattering proﬁles for polymergrafted nanoparticles and polymer nanocomposites at the same relative silica volume fractions and corresponding ﬁts obtained from a polydisperse core−shell model. The curves are arbitrarily shifted for clarity. Transmission electron microscopy images of the (b) 18 kg/mol PNC and (c) 18 kg/mol PGN.
Tg) should be directly associated with the presence of an interfacial polymer in both PNCs and PGNs.31−33 To explore this idea, we use SAXS to determine the peculiarity of the static structure of our materials within the glassy state. Based on our previous results, we found that P2VP/silica nanocomposites exhibit a core−shell structure with a ﬁnite shell thickness.20 In this article, thus we use a polydisperse core−shell model (PCSM) and hard sphere models to ﬁt our SAXS data.34 The properties of the shell in this model are described by the scattering length density, the static shell thickness, and the polydispersity of the shell thickness. These parameters combined provide important information concerning the static length scale, relative density, and the uniformity of the density of the shell. Figure 1a presents the SAXS spectra with the corresponding ﬁts to the best model. The 18 and 140 kg/mol PNC and PGN samples ﬁt to the PCSM model, requiring three structurally diﬀerent regions, core, shell, and polymer matrix. The 6 kg/mol PNC and PGN ﬁt best to the hard sphere model, as no shell was distinguishable. The complete set of the ﬁt parameters can be found in Table S1 in the Supporting Information (SI). For PGNs, the static length scale of the interfacial shell was found to stay relatively constant for the intermediate and high MW with values of 3.4 and 3.8 nm, respectively. Interestingly, the length scale of the interfacial shell was found to be systematically ∼20−35% smaller in the PNCs than in the PGNs. Also, the scattering length density in the interfacial 6844
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Figure 2. (a) Speciﬁc heat capacity from calorimetry illustrating the glass transition of the matrix component for the polymer nanocomposites and polymer-grafted nanoparticles for each molecular weight. (b) Derivative of speciﬁc heat capacity to better demonstrate the broadening of the glass transition in the form of a peak. All curves are arbitrarily vertically shifted for clarity.
Table 1. Summary of Dynamic Information from TMDSC and BDS Combined with Selected Parameters from the Interfacial Model Analysis interfacial layer [nm] type of system
SiO2 (wt %)
SiO2 (vol %)
avg Tg [K] (TMDSC)
avg Tg [K] (VFT)
inter. Tg [K] (VFT)
inter. polymer (vol %)
6 kg/mol, neat 6 kg/mol, PNC 6 kg/mol, PGN 18 kg/mol, neat 18 kg/mol, PNC 18 kg/mol, PGN 140 kg/mol, neat 140 kg/mol, PNC 140 kg/mol, PGN
0 66 65 0 50 50 0 12 12
0 49 48 0 33 33 0 6 6
353.1 365.2 365.6 367.9 371.1 377.0 375.6 376.5 376.5
347.1 362.4 362.3 361.9 367.0 372.4 370.4 370.9 370.7
N/A 362.4 362.3 N/A 371.9 375.7 N/A N/A N/A
N/A ∼100 ∼100 N/A 28 40 N/A N/A N/A
BDS N/A 3.28 ± 3.45 ± N/A 2.81 ± 3.75 ± N/A N/A N/A
0.10a 0.10a 0.11 0.15
SAXS N/A N/A N/A N/A 2.35 ± 0.8 3.4 ± 0.8 N/A 3.1 ± 0.7 3.8 ± 0.9
The dynamic interfacial layer thickness is calculated assuming that 100% of the polymer is interfacial.
region is ∼20% lower than scattering length density of the matrix for the 18 and 140 kg/mol PNCs and PGNs. However, the 6 kg/mol samples are best described by a hard sphere model, which indicates that the matrix surrounding the nanoparticles is homogeneous. This is consistent with a nanocomposite where the entire surrounding matrix is interfacial polymer due to the high silica loading (∼50 vol % silica). The glass transitions of the PGNs and PNCs were measured by TMDSC. Figure 2a shows the speciﬁc heat data after accounting for the silica contributions, as previously explained in references 5 and 20. The glass transition temperatures, Tg (taken as the midpoint of the transition step), for both PNCs and PGNs (Table 1) are found to be dependent on MW (and relative silica concentration). Both PNC and PGN with high MW (140 kg/mol) exhibit a negligible shift in Tg, and the experimental data almost coincide with the data of the neat polymer. This result is likely due to the low silica concentration (6 vol % silica or average interparticle spacing ∼30 nm). The intermediate MW systems show a clear change in Tg, with a larger shift for PGNs (∼9 K) and a smaller shift for PNCs (∼3 K). The low MW (6 kg/mol) samples show a much stronger change in the glass transition than the other MWs. Although the shift in Tg (midpoint) ∼ 13 K appears to be the same for PNC and PGN, a signiﬁcantly broader step is observed in the PGN system. The broadening of the glass transition of the PGN relative to the corresponding PNC can be clearly seen from the derivative analysis of the TMDSC data presented in Figure 2b. The width of the glass transition of the PGN extends
to higher temperatures than for the PNC system. This increase in the breadth of the glass transition indicates two important points: (1) the segmental mobility has a large gradient, and (2) a signiﬁcant amount of the segmental mobility is slower than the average mobility (i.e., midpoint of the transition). Our results for the low MW systems, which have extremely high silica loadings and an estimated interparticle spacing of ∼3.0 nm, are consistent with recent simulations and predictions that segmental dynamics of nanoconﬁned polymers exhibit a large gradient of segmental mobility.3,6,7,11,12,20,35 To better understand the interfacial segmental dynamics and their eﬀect on the overall segmental dynamics in these systems, we employ BDS. The segmental relaxation process appears as a peak in the dielectric loss spectra (Figure 3). We are speciﬁcally interested in the spectral shape of the segmental relaxation for poly(2-vinylpyridine) at temperatures above and close to Tg. The presented dielectric loss data are normalized to the highfrequency tail of the segmental relaxation to compare the shift and overall broadening of the relaxation process at low frequency.20 While the high-frequency tail of the segmental process is slightly aﬀected in both systems, the systematic broadening of the low-frequency part is evident for both PGN and PNC systems. These results corroborate the previously presented calorimetric data. For high MW PNC and PGNs, the width and peak position of the segmental relaxation is essentially unaﬀected, similar to the TMDSC data. However, at intermediate MW (Figure 3b), the dispersion and peak position in the spectra of PNC and PGN are diﬀerent at the same temperature. The corresponding intermediate MW PNC 6845
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Figure 3. Normalized dielectric loss spectra demonstrating the segmental relaxation process for neat polymer, PNCs (red), and PGNs (blue) for each molecular weight: (a) 140 kg/mol, (b) 18 kg/ mol, and (c) 6 kg/mol at select temperatures.
Figure 4. Deconvolution of eﬀective dielectric loss spectra (εeff″(ω)) into individual components: the matrix response (εm″(ω), black squares) and the interfacial layer response (εIL″(ω), blue line) for the intermediate molecular weight (18 kg/mol) of (a) PGN and (b) PNC from the interfacial model analysis (T = 423.15 K). The presented εIL″(ω) component is normalized by the corresponding volume fraction of interfacial polymer. (c) Temperature dependence of the interfacial segmental relaxation times for PGNs (blue), PNCs (red), and bulk-like dynamics (black) with corresponding Vogel−Fulcher−Tammann ﬁts.
system has a peak position similar to that in the pure polymer and a smaller broadening than the PGN. At low MW (Figure 3c), we again observe the most signiﬁcant changes in the dielectric loss spectra, where the entire segmental peak shifts to lower frequencies and clearly broadens for both systems. However, despite the same peak position of the segmental relaxation, there is a clear diﬀerence in the amount of broadening (Figure 3c): the distribution in the PGN system is much wider, similar to calorimetric data in Figure 2b. An indication of interfacial dynamics is the broadening of the segmental relaxation process in the dielectric spectra (Figures 3 and 4) and in the calorimetric data (Figure 2b). Although broadening of the spectra and heterogeneous dynamics in composite materials has been reported and associated with the formation of the interface, traditional analysis is often limited to the comparison of the average values for Tg (midpoint of transition) and relaxation time (peak position), resulting in averaged information about the properties of the material. However, the properties of the interfacial region itself are extremely important as they impact the macroscopic properties of PNCs that are coupled to segmental dynamics, such as ion transport (e.g., polymer electrolyte applications) or smallmolecule diﬀusion (e.g., gas separation).36−38 To reveal the properties of the interfacial layer, we utilize the interfacial model, which describes an eﬀective medium approximation for the dielectric response (εeff″(ω)) of a heterogeneous system, accounting for the dielectric response of the particle (εNP″(ω)), interfacial layer (εIL″(ω)), and matrix (εm″(ω)). The speciﬁc details of this analysis are included in the SI and are also welldescribed in our earlier publications.19,39 Using the interfacial model, we can determine two critical parameters: (1) the time scale of the interfacial segmental dynamics and (2) the volume fractions of polymer populations that are bulk-like or interfacial. Here, we analyze the data for these PGNs and PNCs at the intermediate MW because they demonstrate the most
pronounced change in the dispersion of the segmental relaxation. The interfacial model analysis (Figure 4) separates the contributions of the interfacial and bulk-like segmental dynamics and provides an estimate for the respective volume fractions and characteristic relaxation times. For the intermediate MW systems, the characteristic relaxation times of the interfacial layer are determined by ﬁtting the obtained εIL″(ω) with a single Havriliak−Negami function.40 At high temperatures, the segmental relaxation times of the interfacial and bulk-like dynamics converge (Figure 4c), indicating more homogeneous dynamics. However, even at high temperatures, the segmental dynamics in the PGN remain slower than in the neat polymer. As the temperature approaches Tg (τα = 100 s, as indicated by the horizontal dotted line), the slowing down of the interfacial dynamics relative to the neat polymer increases strongly, especially for the PGN system. The dynamic glass transition temperatures, Tg, can be determined by ﬁtting the relaxation times to the Vogel−Fulcher−Tammann (VFT) equation and extrapolating to the temperature at which τα = 100 s (Table 1). A diﬀerent behavior is observed for PNC and PGN for the lowest MW samples (Figure 5). In that case, the characteristic relaxation times are extracted by ﬁtting the entire εeff″(ω) to a single Havriliak−Negami function because essentially the entire sample is an interfacial region at this high loading. Both PNC and PGN exhibit a similar time scale of segmental dynamics at T close to Tg (consistent with calorimetric data), while the PGN shows stronger deviations from the behavior of neat polymer than PNC at higher T. The observed suppression of interfacial dynamics is in line with several molecular dynamic simulations2,3,6,7 (which spatially map the distribution of relaxation times) and qualitatively agrees with the expectation that covalently bonded chains experience a larger suppression of segmental dynamics, depending also on the grafting density6 and molecular weight.41 6846
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interfacial polymer. Therefore, the interfacial dynamics contribution is too weak to be detected by our experimental methods, and we exclude it from the rest of the discussion about dynamics. However, the intermediate MW systems show rather large changes in the segmental relaxation process and also a clear diﬀerence in the time scale of the interfacial dynamics between PGNs and PNCs (Figure 4c). One of the main diﬀerences in these dynamics can be related to the volume fraction of the interfacial polymer, which is ∼0.6 in the PGN system while is only ∼0.4 in the PNC system. Most interestingly, analysis of the low MW systems, which have extremely high nanoparticles loading (∼50 vol % silica), reveals that the interfacial region occupies the entire polymer matrix, regardless of physical adsorption or covalent bonding, as reﬂected in the average segmental dynamics (Figure 3c and Figure 5). Additionally, the length scale of the interfacial region can be estimated from the volume fraction of the interfacial polymer by assuming a core−shell morphology, as described previously.20 The dynamic length scale of the interfacial layer estimated from the BDS measurements is also compared to the static shell thickness estimated from SAXS (inset Figure 6a). For PGNs, the interfacial layer thickness seems to slightly decrease with decreasing MW, while no systematic changes can be found in PNCs as a function of MW. It seems that for all systems exhibiting a distinct interfacial layer, the static interfacial layer thickness is larger for PGNs than for PNCs, although the diﬀerence is within the experimental error-bars (inset Figure 6a). Our earlier studies on PNCs with ﬁxed NP loading19 revealed a decrease in the interfacial layer thickness and signiﬁcant decrease of the averaged Tg shift with increasing molecular weight (see also Figure S7 in SI). However, it is diﬃcult to identify a certain trend in the current data set because we cannot detect the static length at low molecular weight or the dynamic length at high molecular weight. Our analysis also reveals that the dynamic length scale of the interfacial layer increases slightly upon cooling (Figure S6 in SI), consistent with our earlier observations in model glycerol/ SiO2 nanocomposites.39 The observed increase in the dynamic length scale of the interfacial layer upon approaching Tg might be related to an increase in cooperativity of the segmental dynamics.2,3 However, this discussion is out of scope of the current study. Figure 6b presents the relative change in Tg (with respect to neat polymers) estimated using three techniques: (1) the average Tg from TMDSC (midpoint of transition), (2) the average Tg from BDS (using a single Havriliak−Negami relaxation process to describe the εeff″(ω) spectra and VFT extrapolation), and (3) the interfacial Tg from the interfacial model analysis (using a single Havriliak−Negami relaxation process to describe the εIL″(ω) spectra and VFT extrapolation). From Figure 6b, it becomes clear that the low MW PNCs and PGNs experience similar interfacial dynamics resulting in the same average Tg from TMDSC and BDS. It is diﬃcult to draw any conclusions from the high molecular weight systems due to the relatively low fraction of the interfacial region at such a small content of silica nanoparticles. However, the intermediate MW PNC and PGN show a clear diﬀerence in the average Tg from both techniques, as well as a large diﬀerence in the interfacial Tg. Following our earlier studies, we relate these observations to the diﬀerences in chain conformations within the interfacial layer.19
Figure 5. Temperature dependence of the segmental relaxation times for low molecular weight (6 kg/mol) PGNs (blue), PNCs (red), and bulk-like dynamics (black) with the corresponding VFT ﬁts (lines).
To rationalize these results, we must ﬁrst understand how much of the polymer matrix is actually inﬂuenced by the nanoparticles (i.e., the volume fraction of the interfacial regions). Figure 6a presents the normalized volume fraction of the interfacial polymer (obtained from the interfacial model) for both PGNs and PNCs at the temperatures presented in Figure 3. Also, the temperature dependences of these values are included in the SI. At high MW, there is a small amount of nanoparticles (6 vol % silica) and, as a result, even less
Figure 6. (a) Relative volume fraction of interfacial polymer with respect to the total polymer for both systems as a function of molecular weight. Inset: Comparison of the static and dynamic thickness of the interfacial region estimated using BDS (full symbols) and SAXS (empty symbols) as a function of MW. (b) Change in the average calorimetric glass transition temperature, Tg, for PGNs and PNCs from TMDSC as well as the average and interfacial Tg from BDS by evaluating the VFT function for each molecular weight as a function of silica concentration. The boxes indicate data at the same respective MWs, and the dashed lines are guides for the eye. 6847
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expected for adsorbed polymers. The polymer−nanoparticle interactions were modeled via a step function which only acts on the exterior of the nanoparticles. The range of the interaction potential was kept ﬁxed at 1.7 nm, representing the rough native oxide layer present on the silica nanoparticles.51 The strength of the interaction potential was chosen to match the line averaged density proﬁle for PNCs (between the nanoparticles) obtained from simulations with the experimentally obtained spatially averaged density proﬁle (∼1.0 g/cm3). As shown in Table 2, h/2Rg systematically increases from ∼1 at low MW to nearly 1.5 at high MW. This suggests that the polymer chains exhibit highly stretched conformations at the interface and that the degree of stretching increases with MW for the PGN systems. To understand chain stretching for the case of physical adsorption, we use SCFT to calculate the distribution of chain conformations for the PNC systems. As shown in Figure 7, the distribution of chain conformations was calculated for our exact systems and provides quantitative information concerning the spatial distribution of polymer segments near the spherical nanoparticle interfaces. To quantify the degree of chain stretching in PNCs, we analyzed the distribution of tails, loops, and trains and the length scale that they persisted into the free polymer population far away from the surface of the nanoparticles, dcomp. At low MW, we ﬁnd that the maximum populations of tails and loops are spatially distributed further away from the NP surface than the length scale of the unperturbed Rg (Figure 7a). For intermediate MW, we observe fewer tail conformations with a more signiﬁcant population of loops and trains, especially near the NP interface (Figure 7b). The maximum population of tail conformations are now found at distances approaching the Rg, and only a small fraction of the tails propagate beyond Rg. Additionally, the interfacial region up to 2.5 nm largely consists only of segments that are in direct contact with the nanoparticle (Figure 7b). For high MW, the majority of the chain conformations actually consist of tails that persist up to 12 nm (approximately Rg) away from the NP surface, where they mix with the free polymer chains (Figure 7c). The spatial distributions of the tails, loops, and trains depend on interplay of adsorption energy and chain conformational entropy. For shorter chains, the conformational entropy is weaker compared to the interaction energy between the polymer chains and the nanoparticle surfaces. This, in turn, leads to more loops and trains in comparison with the number of tails. In contrast, the conformational entropy of the polymer chains dominates over the chain−nanoparticle interactions for the longest chains studied in this work, which results in a larger number of tail conformations than loop or train conformations. To quantify the degree of chain stretching for all PNCs, we determine the ratio of the distance that the adsorbed polymer segments persist into the free polymer segments, dcomp, and the unperturbed Rg, a parameter that is similar to h/2Rg and is presented in Table 2. The stretching factor, dcomp/2Rg, of the PNCs decreases with increasing MW, behavior that is opposite to that of the PGN systems. Interestingly, the stretching factor for both systems converge at low MW (h/2Rg ≈ dcomp/2Rg ≈ 1.0), suggesting that the same degree of stretching exists in these two systems. These ﬁndings explain why the low MW PNC and PGN systems exhibit the largest changes in segmental dynamics. For the intermediate MW, the stretching factors also agree with the corresponding diﬀerences in the interfacial dynamics between PNC and PGN: the more strongly stretched PGN system (h/
In general, it is diﬃcult to directly assess the conformations of polymers in PNCs with physically adsorbed chains. However, the conformations of covalently attached chains are very well studied experimentally42 and can be described theoretically.43−50 Polymers attached to a surface experience diﬀerent conformations ranging from mushroom-like structures at low grafting density, where neighboring sites do not overlap, to highly stretched conformations at high grafting density. In general, chain stretching can be described by the dimensionless parameter σ∥ = h/2Rg, where h is the brush height for a given interparticle spacing (IPS) and Rg is the radius of gyration of a Gaussian chain in the polymer melt. By this deﬁnition, the interfacial chain conformations are considered Gaussian when σ∥ is much less than one and stretched when σ∥ is one or greater. The speciﬁc values of σ∥ for the experimental systems studied in this work are listed in Table 2, which are based on Table 2. Calculated Stretching Parameters for Both Polymer Nanocomposites and Polymer-Grafted Nanoparticles MW
6 kg/mol 18 kg/mol 140 kg/mol
3.3 7.0 30
4.64 7.9 22.2
4.7 9.0 32.5
4.6 6.4 12.3
1.01 1.13 1.46
0.99 0.81 0.55
the results obtained from the SCFT calculations in a twodimensional polar coordinate system (see details of the calculations in the SI). The SCFT calculations were done for a compressible system. In addition, the attractive interactions between the polymer segments and the nanoparticle surface were included for both the PGN and PNC systems. Due to interplay of the depletion region next to the nanoparticle surface (resulting from the Dirichlet boundary conditions for propagators in the SCFT) and attractive interactions between the polymer segments and the nanoparticle, nonmonotonic density proﬁles (Figure 7) are obtained for these systems, as
Figure 7. Self-consistent ﬁeld theory calculations for all the polymer nanocomposites. Here, we illustrate the diﬀerence in the adsorbed chain conformations near the nanoparticle’s surface by comparing the distribution of adsorbed segments and the length scale that they persist into the free polymer (dashed vertical black lines, dcomp) relative to the matrix’s unperturbed radius of gyration (dashed vertical pink lines) to quantify the degree of chain stretching. 6848
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ACS Nano 2Rg ≥ 1.0) exhibits a larger suppression of interfacial dynamics, whereas the less stretched PNC system (dcomp/2Rg < 1.0) exhibits a less pronounced eﬀect. The results above suggest that the mechanism responsible for the suppression of the segmental dynamics at the interface is directly correlated with the degree of stretching of the polymer chains at the polymer−NP interface (Figure 8) and might be
CONCLUSIONS The interfacial structure and dynamics of P2VP/silica nanocomposites and matrix-free P2VP grafted from silica nanoparticles are studied under identical conditions. We utilize an interfacial layer model to unravel the time scale of the interfacial segmental dynamics from the dielectric relaxation spectra. The high MW systems do not show signiﬁcant variations due to relatively low volume fraction of nanoparticles and interfacial region. However, a strong gradient of segmental mobility in the interfacial region was revealed by TMDSC and BDS techniques for both low and intermediate MW PNCs and PGNs. At intermediate MW, our analysis reveals a signiﬁcant inﬂuence of the chain attachment mechanism, such as physical adsorption of the chains for PNC versus covalent grafting of the chains for PGNs, on segmental dynamics and Tg of the material. Interestingly, for low MW PNCs, the dynamic properties were found to be similar to low MW PGNs, indicating that these systems experience a similar eﬀect at the interfaceindependent of the method of their preparation. In lieu of these results, we propose that the chain stretching at the interface, rather than the polymer−nanoparticle interaction strength, plays a critical role in suppression of segmental dynamics, which ultimately determines the magnitude of the change in the average glass transition temperature in these heterogeneous materials. In other words, the intermediate MW PGNs exhibit a higher degree of chain stretching at the interface and therefore are accompanied by a more dramatic change in Tg when compared to their PNC counterparts, whereas the low MW systems share a similar degree of chain stretching and also exhibit a similar change in dynamics. Overall, these results demonstrate that polymer chain conformations in the interfacial region can strongly inﬂuence the interfacial segmental dynamics and provide physical insight concerning dynamic interfaces in polymer nanocomposites and other nanostructured polymer systems, thereby suggesting a way to rationally tailor polymer dynamics on the nanoscale.
Figure 8. Cartoon based on the SCFT results that illustrates how chain stretching can cause the suppression of segmental mobility at the interface in polymer-grafted nanoparticles (a,c) and polymer nanocomposites (b,d) at diﬀerent molecular weights.
less inﬂuenced by how the chain is attached to the NP, for example, physical adsorption or covalent bonding. This phenomenon is likely due to the diﬀerence in the chain crowding at the interface, where chain crowding imposes stronger steric hindrances at the NP surfaces, an eﬀect that should quickly dissipate with increasing distance from the NP surface (unlike planar systems). Therefore, the curvature of the NP should play a signiﬁcant role in this eﬀect, and one should expect larger changes for larger NPs or planar surfaces. This was independently reported by Gong et al.,52 where the interfacial thickness was found to increase with increasing NP sizeconsistent with the concept of chain crowding and chain stretching. However, we speculate that there is a certain trade oﬀ with NP size because the lower surface-to-volume ratio of larger NPs would reduce the volume fraction of the interfacial region in the nanocomposite with the same loading. From a fundamental perspective, these results qualitatively agree with previous predictions by Oyerokun and Schweizer30 for the dynamics of polymer systems with anisotropic conformations (thin ﬁlms, planar grafted brushes) which predicted that segmental stretching increases the correlation between segmental motions (intrachain) and increases the local Tg at the interface. We speculate that this eﬀect will be more pronounced for rigid polymer backbones, as demonstrated in recent molecular dynamic simulations.53 It should be noted that the results presented in this article are speciﬁc to polymer−nanoparticle systems with strong attractive interactions. The behavior might be diﬀerent if the time scale of chain adsorption (to the nanoparticles) is comparable to the segmental relaxation time. In general, we would expect that PNC systems and PGNs may exhibit signiﬁcant diﬀerences in the dynamics depending on nanoparticle loading, molecular weight, or grafting density, especially in the case of unfavorable (repulsive) polymer−nanoparticle interactions. That does not necessarily imply that the underlying mechanism is diﬀerent, but rather that the magnitude of the eﬀect will be diﬀerent and that covalent bonding may provide a route to modify the segmental dynamics of polymer−nanoparticle systems with naturally weak polymer−nanoparticle interactions.
METHODS Materials. 2-Vinylpyridine (2VP, 97%, Aldrich) was deinhibited via passing through an activated alumina column. 4-Cyano-4(phenylcarbonothioylthio)pentanoic acid N-succinimidyl ester (RAFT agent, Aldrich), hexanes (BDH), and tetrahydrofuran (THF, BDH) were used as received. 2,2-Azobis(isobutyronitrile) (AIBN, 99%, Aldrich) was puriﬁed by recrystallization prior to use. Synthesis of Silica Nanoparticles and Attachment of RAFT Agent. Silica nanoparticles with the diameter of 20−25 nm were synthesized following the modiﬁed Stöber method. The details of the synthesis are described elsewhere.54,55 As-prepared silica was modiﬁed with 3-aminopropyldimethylethoxysilane. For that, 100 mL of silica nanoparticles (15 mg/mL) in ethanol was rotary evaporated and redispersed in 100 mL of dry THF; 1 mL of silane was added to the silica nanoparticles in THF, and the mixture was reﬂuxed for 24 h. The particles were centrifuged three times to remove any unreacted silane. The washed particles were then redispersed in 50 mL of dry THF, and 0.1 g of RAFT agent was added to the solution. The resulting mixture was stirred at room temperature for 12 h. Centrifugation was used to separate unreacted RAFT agent from RAFT-agent-functionalized silica particles. Synthesis of Poly(2-vinylpyridine) on Silica Nanoparticles via RAFT Polymerization. 2-Vinylpyridine (19.92 g, 20.43 mL) and AIBN (4.98 × 10−5 mol, 8.18 mg) were put into RAFT-agentfunctionalized silica particle THF solution (48 mL). Argon was bubbled through the solution for 10 min. The reaction proceeded for 2 h 15 min, 6 h, and 18 h at 65 °C. The resulting poly(2-vinylpyridine) (P2VP) grafted from silica NPs was precipitated into hexanes, 6849
DOI: 10.1021/acsnano.6b02501 ACS Nano 2016, 10, 6843−6852
ACS Nano centrifuged, andwashed with THF/hexane mixtures three times. The supernatant was collected, and the unattached P2VP was precipitated into hexanes. The precipitated particles and isolated free P2VP were dried at 40 °C in vacuum for 24 h. Molecular weight characteristics of the unattached polymers were obtained from size-exclusion chromatography (SEC) equipped with both light scattering and refractive index detectors. The measured molecular weight of the free polymers was used to determine the molecular weight of the grafted P2VP. The free polymer was then saved and used for the preparation of the polymer nanocomposites. Our SEC system consisted of a Waters 2695 Alliance HPLC pump equipped with a degasser and an autosampler, 3X Polymer Laboratories mixed-C ultrapolystryagel columns in a thermostated compartment, a Wyatt miniDAWN 3angle ambient light scattering detector, and a Waters 2414 refractive index detector; N,N-dimethylformamide was used as the eluent at a 0.5 mL/min ﬂow rate with the columns at 60 °C. MW values were preferentially taken from light-scattering-based calculations using Wyatt Astra software. Conventional calibration MW and polydispersity index values were also taken according to the RI detector results and Waters Empower software; third-order polynomial calibrations were made using either P2VP standards or polystyrene standards, and these standards were also passed the same day as unknown analytes to verify calibration durability. Thermogravimetric analysis was used to determine the weight fraction of silica and grafted P2VP. This information together with the molecular weight obtained from GPC was used to estimate grafting density of P2VP, which was found equal to 0.3 chains/nm2 for all the polymer-grafted nanoparticle samples. Preparation of Polymer Nanocomposites. The free (unattached) P2VP chains collected from the grafted nanoparticle synthesis were used to prepare composites with identical silica nanoparticles (with unmodiﬁed surface functionality) with the molecular weights identical to the grafted samples. The polymer and nanoparticles were mixed in ethanol; the solvent was evaporated during stirring, and samples were dried in the Teﬂon plates at 100 °C for 48 h. The amount of silica nanoparticles in composite materials was added at the identical weight fraction of P2VP than that of the covalently grafted systems and conﬁrmed by thermogravimetric analysis. For TEM measurements, the samples were embedded in epoxy, cured overnight, and subsequently microtomed with a Leica Ultra microtome EM UC7 using a glass knife at ambient conditions, with resulting specimens having thicknesses of ∼100 nm. The TEM measurements were performed on Zeiss Libra 200 HT FE MC with an operating voltage of 200 kV and an emission current of 230 mA. Small-angle X-ray scattering measurements were performed using a high ﬂux small-angle X-ray scattering instrument (SAXSess mc2, Antom Paar, Austria) equipped with a Kratky block-collimation system and a sealed-tube Cu Kα X-ray generator with the wavelength of λ = 0.154 nm, operating at tension = 40 kV and tube current = 50 mA at the Center for Nanophase Materials Science at Oak Ridge National Laboratory. The scattered X-ray intensities were detected by a 2D charge-coupled detector with a spatial resolution of 24 × 24 μm2 per pixel at a sample-to-detector distance of 30.9 cm. The scattering intensities measured as a function of half the scattering angle, θ, were ﬁrst corrected for the absorption of the X-rays by the sample followed by further correction for the scattering of the empty cell and transformed to a plot of scattering intensity versus momentum transfer vector, Q (Q = (4π sin θ)/λ). The corrected data were normalized with respect to the scattering of a water sample that is measured at exactly the same conditions as the samples. Temperature-modulated diﬀerential scanning calorimetry measurements were carried out on a Q2000 (TA Instruments) and were calibrated by indium and sapphire standards using standard aluminum pans. The samples were equilibrated at 473.15 K for 30 min and then cooled to 300 at 5 K/min with a modulation of ±1 K/min. All TMDSC measurements were repeated three times. BDS measurements were performed using a Novocontrol Alpha analyzer, and the sample temperature was controlled by a Quatro Cryosystem with a stability of ±0.1 K. Prior to measurements, the samples were annealed at 473.15 K in the dielectric sample chamber,
and the response was measured as a function of time until the real and imaginary permittivity became constant for any given frequency. The measurements were performed in the frequency range of 10−1 to 107 Hz, and all samples were measured twice to ensure data reproducibility.
ASSOCIATED CONTENT S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.6b02501. A detailed explanation of the small-angle X-ray scattering analysis, the interlayer model analysis used for the dielectric spectroscopy data, and self-consistent ﬁeld theory calculations (PDF)
AUTHOR INFORMATION Corresponding Authors
*E-mail: [email protected]
*E-mail: [email protected]
The authors declare no competing ﬁnancial interest.
ACKNOWLEDGMENTS We thank Ken Schweizer and Eileen Buenning for many useful and insightful discussions. This work was supported by the U.S. Department of Energy, Oﬃce of Science, Basic Energy Sciences, Materials Sciences and Engineering Division. We also acknowledge the support of the Scientiﬁc User Facilities Division, Oﬃce of Basic Energy Sciences, U.S. Department of Energy, who sponsors the Center for Nanophase Materials Sciences (CNMS) at Oak Ridge National Laboratory. N.S. thanks the NSF-REU summer program (CHE-1262767) at the University of Tennessee for the opportunity to work on this project. REFERENCES (1) 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. (2) Starr, F. W.; Douglas, J. F. Modifying Fragility and Collective Motion in Polymer Melts with Nanoparticles. Phys. Rev. Lett. 2011, 106, 115702. (3) Pazmiño Betancourt, B. A.; Douglas, J. F.; Starr, F. W. Fragility and Cooperative Motion in a Glass-Forming Polymer−nanoparticle Composite. Soft Matter 2013, 9, 241. (4) Starr, F. W.; Schrøder, T. B.; Glotzer, S. C. Molecular Dynamics Simulation of a Polymer Melt with a Nanoscopic Particle. Macromolecules 2002, 35, 4481−4492. (5) Sargsyan, A.; Tonoyan, A.; Davtyan, S.; Schick, C. The Amount of Immobilized Polymer in PMMA SiO2 Nanocomposites Determined from Calorimetric Data. Eur. Polym. J. 2007, 43, 3113−3127. (6) Ghanbari, A.; Rahimi, M.; Dehghany, J. Influence of Surface Grafted Polymers on the Polymer Dynamics in a Silica-Polystyrene Nanocomposite: A Coarse-Grained Molecular Dynamics Investigation Influence of Surface Grafted Polymers on the Polymer Dynamics in a Silica-Polystyrene Nanocomposite. C. J. Phys. Chem. C 2013, 117, 25069. (7) Ghanbari, A.; Ndoro, T. V. M; Leroy, F.; Rahimi, M.; Böhm, M. C.; Müller-Plathe, F. Interphase Structure in Silica-Polystyrene Nanocomposites: A Coarse-Grained Molecular Dynamics Study. Macromolecules 2012, 45, 572−584. (8) Robertson, C. G.; Hogan, T. E.; Rackaitis, M.; Puskas, J. E.; Wang, X. Effect of Nanoscale Confinement on Glass Transition of 6850
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