Polymethacrylates with Polyhedral Oligomeric Silsesquioxane (POSS

Article pubs.acs.org/Macromolecules

Polymethacrylates with Polyhedral Oligomeric Silsesquioxane (POSS) Moieties: Influence of Spacer Length on Packing, Thermodynamics, and Dynamics Stelios Alexandris,† Adrian Franczyk,‡,§ George Papamokos,† Bogdan Marciniec,§ Krzysztof Matyjaszewski,‡ Kaloian Koynov,∥ Markus Mezger,∥ and George Floudas*,† †

Department of Physics, University of Ioannina, 45110 Ioannina, Greece Department of Chemistry, Carnegie Mellon University, 4400 Fifth Avenue, Pittsburgh, Pennsylvania 15213, United States § Center for Advanced Technologies, Adam Mickiewicz University in Poznan, Umultowska 89c, 61-614 Poznan, Poland ∥ Max-Planck Institute for Polymer Research, 55128 Mainz, Germany ‡

S Supporting Information *

ABSTRACT: Polymethacrylates with polyhedral oligomeric silsesquioxane (POSS) moieties (poly(POSS-MA)s) with flexible spacers between the POSS cages and the methacrylate group have distinctly different properties from their linear counterpart, i.e., PMMA. POSS cages modify interchain correlations and result in multiple dynamic processes that reflect the cooperative relaxations of both the pendant POSS units and ester dipoles and the polymer backbone. As a result, the freezing of the backbone dynamics is shifted to lower temperatures, and the nanocomposites appear softer than linear PMMA chains of similar degrees of polymerization. POSS cages can be employed as nanometer size blocks that, depending on the polymer backbone and the spacer, can impart mobility and control over the mechanical properties of nanocomposites. also some reports either of unaltered7 or reduced Tg in poly(acetoxystyrene)-POSS copolymers.6 Recently, a new polyhedral oligomeric silsesquioxane methacrylate monomer ((i-Bu)7POSS-OSiMe2(CH2)3-MA) with seven small isobutyl substituents and a flexible spacer between the POSS cage and the methacrylate group was synthesized.8 Atom transfer radical polymerization (ATRP)9 of this monomer resulted in polymethacrylates (poly(POSSMA)s) of very high molecular weight (Mw > 105). Previous attempts of preparation of high molecular weight poly(POSSMA) with larger substituents resulted in only low molecular weight polymers.10 Higher initial monomer concentration also helps to increase polymer yields and molecular weight.8,11 The flexible spacer plays an essential role as it relaxes the strain generated by the sterically bulky cage through the longer Si− O−Si bonds and enhances accessibility of the end group for the monomer and the catalyst in the ATRP process. Furthermore, the flexible spacer can also facilitate a better packing of POSS units. However, the effect of related changes in packing on the thermodynamics, the local dynamics, and eventually the

1. INTRODUCTION Hybrid polymer/ceramic building blocks combine properties of organic polymers, such as processability and toughness, with properties of inorganic compounds (i.e., thermal and chemical stability). Recently, a polyhedral oligomeric silsesquioxane (POSS) moiety as nanometer-size building block has been successfully incorporated into a range of polymeric materials.1 Cubic siloxane cages are unique in that they are physically large (i.e., the POSS units can be considered as giant atoms)2 and yet composed of a robust silicon−oxygen framework that can be easily functionalized at the corners of the cage with different organic substituents. Furthermore, POSS reagents can be incorporated into polymer systems using established polymerization protocols. Incorporation of POSS segments to linear polymers may result in alteration of the glass temperature (Tg), decomposition time and temperature, increased oxygen permeability, and reduced flammability as well as in modified mechanical properties in comparison to non-POSS-containing polymers.1 Moreover, these properties can be rationally tailored by controlling the amount of POSS moieties. With respect to Tg, a hindered chain relaxation was reported when POSS units were chemically incorporated directly into the polymer chain. This was the case of norbonyl-POSS hybrid copolymers,3 styryl-POSS copolymers,4 monofunctional epoxy-POSS,5 and poly(vinylpyrolidone)-POSS copolymers.6 However, there are © XXXX American Chemical Society

Received: March 30, 2015 Revised: May 3, 2015

A

DOI: 10.1021/acs.macromol.5b00663 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules

increment of the samples was determined assuming 100% mass recovery. All data were recorded and evaluated using WinGPC UniChrom software from PSS. Nuclear Magnetic Resonance Spectrometry (NMR). 1H, 13C, and 29 Si NMR spectra were recorded on a Varian Gemini 300 VT spectrometer and Varian Mercury 300 VT in CDCl3. Thermogravimetric Analysis (TGA). TGA was performed using a TGA Q50 V20.13 Build 39, and the data were recorded over a temperature range of 298−873 K under a nitrogen atmosphere at a heating rate of 10 K/min. Differential Scanning Calorimetry (DSC). A Q2000 (TA Instruments) was used for thermal analysis with a cooling/heating rate of 10 K/min at a temperature range from 153 to 413 K. The instrument was calibrated for best performance on the specific temperature range and heating/cooling rate. The calibration sequence included a baseline calibration for the determination of the time constants and capacitances of the sample and reference sensor using a sapphire standard, an enthalpy and temperature calibration for the correction of thermal resistance using indium as standard (ΔH = 28.71 J/g, Tm = 428.8 K), and a heat capacity calibration with sapphire standard. Temperature-modulated DSC (TMDSC) was made with an amplitude of 1 V and for periods in the range from 20 to 150 s. X-ray Scattering. X-ray diffraction (XRD) on monomer and polymer samples was recorded with Cu Kα radiation (λ = 1.54 Å) using a powder diffractometer in Bragg−Brentano geometry with graphite monochromator. Temperature-dependent XRD was measured in transmission geometry using Cu Kα radiation (Rigaku MicroMax 007 X-ray generator, Osmic Confocal Max-Flux curved multilayer optics). The 1 mm thick polymer sample was contained in a temperature-controlled stainless steel cell with 300 μm thick diamond windows. Diffraction patterns were obtained at a sample−detector distance of 325 mm by radial averaging the 2D data recorded on a Mar345 image plate. Dielectric Spectroscopy (DS). The sample cell consisted of two electrodes, 20 mm in diameter and a thickness of 50 μm. Dielectric measurements were made at different temperatures in the range 253.15−423.15 K, at atmospheric pressure, and for frequencies in the range from 1 × 10−2 to 1 × 106 Hz using a Novocontrol Alpha frequency analyzer with an active sample head. The complex dielectric permittivity ε* = ε′ − iε″, where ε′ is the real and ε″ is the imaginary part, is a function of frequency ω, temperature T, and in general pressure P, ε* = ε*(ω,T,P).12 In the analysis of the DS spectra we have used the empirical equation of Havriliak and Negami (HN)13

viscoelastic properties of nanocomposites needs to be determined. Herein, we employ a series of poly(POSS-MA)s and investigate the effect of the spacer length and polymer molecular weight on the self-assembly as well as the thermodynamic and dynamic properties of the nanocomposites. We employ X-ray diffraction, differential scanning calorimetry, dielectric spectroscopy (DS), and rheology. Density functional theory calculations are employed in extracting the monomer dipole moment in the gas phase in conjunction to the DS study. We established that the presence of pendant POSS cages and the spacer drastically modify the thermodynamic properties and interchain correlations. This, in turn, results in a reduction of Tg and drastically different mechanical properties in comparison to linear poly(methyl methacrylate) (PMMA) chains. We show that POSS moieties impart viscoelasticity in the absence of an elastic plateau.

2. EXPERIMENTAL SECTION 2.1. Polymerization Methods. Polymerization of monomers C3, OSiMe2C3, and OSiMe2C11 by ATRP was performed by methods previously described in the literature.8 Synthesis of Polymers P1−P4 and P6. To a glass flask with a magnetic stir bar, (i-Bu7)POSS-MA, PMDETA, and EBiB were added. In a second flask, CuBr2 (0.012 g, 0.0054 mmol) with PMDETA (0.0112 mL, 0.0054 mmol) was dissolved in DMSO (0.2 mL), and an appropriate volume of this solution was added via syringe to the reaction flask. The reaction mixture was degassed by at least three freeze−pump−thaw cycles and filled with argon (Ar) again. With positive pressure of Ar, CuBr was added to the flask. The flask was evacuated and refilled with Ar and placed in a 323 K oil bath. Samples were taken periodically to measure conversion via 1H NMR and number-average molecular weights via GPC. Synthesis of Polymer P5. A reaction mixture prepared according to the procedure for P1−P4 and P6 was placed in a 323 K oil bath. In a second flask, monomer was dissolved in toluene (0.5 mL), and the air was removed from the solution by at least three freeze−pump−thaw cycles and refilling atmosphere with Ar. The flask was placed in a warm water bath. The clear solution formed by this procedure was added dropwise via syringe to the reaction mixture. Isolation of Polymers P1−P6. The solutions of polymers (mixture of toluene and hexane) were filtrated through the neutral alumina to remove catalyst. Subsequently solutions were concentrated and acetone was added. Resulted “foams” were dried under vacuum. Polymers were washed with acetone to remove the contamination of monomer and dried under vacuum. If the form of polymer was different than powder, the flask with polymer was cooled down in liquid nitrogen, and the resulted solid lumps were crushed and dried under vacuum. 2.2. Characterization Methods. Gel Permeation Chromatography (GPC). MW and MWD of the formed polymers were measured by gel permeation chromatography using Polymer Standards Services (PSS) columns (guard, 105, 103, and 102 Å), with THF eluent at 308 K, flow rate 1.00 mL/min, and differential refractive index (RI) detector (Waters, 2410). Diphenyl ether was used as the internal standard to correct for any fluctuation of the THF flow rate. The number-average MW (Mn,GPC) and MWD (Mw/Mn) were determined with a calibration based on linear poly(methyl methacrylate) (PMMA) standards using WinGPC 6.0 software from PSS. Multiangle Laser Light Scattering Gel Chromatography (MALLS). The measurements were performed on an Agilent 1260 Infinity system coupled to a DAWN-EOS laser light scattering detector (Wyatt Technologies) equipped with a He−Ne laser operating at the wavelength of 685 nm and an ERC refractive index detector. Chromatographic separation was performed by using a combination of three PSS SDV columns, 10 μm, 8 × 300 mm, with porosities: 106, 105, and 500 Å. All experiments were performed in THF at a flow rate of 1 mL/min and at temperature of 303 K. The refractive index

* (ω , T ) = ε∞(T ) + εHN

σ0(T ) Δε(T ) m n + [1 + (iωτHN(T )) ] iεf ω

(1)

where τHN(T,P) is the characteristic relaxation time, Δε(T,P) = ε0(T,P) − ε∞(T,P) is the relaxation strength of the process under investigation, m and n (with limits 0 < m, mn ≤ 1) describe respectively the symmetrical and unsymmetrical broadening of the distribution of relaxation times, σ0 is the dc conductivity, and εf is the permittivity of the free space. In the fitting procedure, we have used the ε″ values at every temperature, and in some cases the ε′ data were also used as a consistency check. From, τHN the relaxation time at maximum loss, τmax, is obtained analytically following

⎛ πm ⎞ 1/ m⎛ πmn ⎞ τmax = τHN sin−1/ m⎜ ⎟ ⎟ sin ⎜ ⎝ 2(1 + n) ⎠ ⎝ 2(1 + n) ⎠

(2)

In the temperature range where two relaxation processes contribute to ε* there are two ways of representing the data. The first one, followed here, is based in a summation of two HN functions and assumes statistical independence in the frequency domain. The second one, proposed by Williams and Watts, is a molecular theory for the dipole moment time-correlation function Cμ(t) (also known as “Williams ansatz”14). An alternative representation of the dielectric data is through the inverse of the dielectric permittivity ε*(ω), i.e., the electric modulus, which is related to the dielectric permittivity through B


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Macromolecules Chart 1. Structures and Thermal Properties of (i-Bu)7POSS-Y-MA Monomers

Chart 2. Solubility of (i-Bu)7POSS-Y-MA Monomers in Toluene at Room Temperature

M *(ω) =

1 = M′ + iM″ ε*(ω)

of 0.14 D (see Table S1, Supporting Information) from larger and sufficient basis sets for the accurate calculation of electric properties.17 For (i-Bu)7POSS(CH2)3-MA and (i-Bu)7POSS-OSiMe2(CH2)11-MA the antiperiplanar conformation of the MA-(CH2)3 and MA-(CH2)11 parts was adopted for the initial structures and remained after optimization. Local minima were confirmed by subsequent normalmode analysis which resulted in positive frequencies for all modes. The calculations were performed employing the Gaussian 03 software package18 while for the visualization of the results the VMD software was used.19 Cartesian coordinates of the optimized structures are available in the Supporting Information.

(3)

In eq 3, M′ and M″ are the real and imaginary parts of the electric modulus, respectively. The electric modulus representation, M*(ω), is an absolute necessity in cases where a comparison with rheology data is needed as in the current work. The relaxation times obtained from the electric modulus (τM*) and the complex permittivity (τε*) representations scale as15 −1/ m ⎛ Δε ⎞ τM ″ ∼ τε ″⎜1 + ⎟ ε∞ ⎠ ⎝

(4)

3. RESULTS AND DISCUSSION Monomers Characterization and Polymerization Results. (i-Bu)7POSS-(CH2)3MA (C3), provided by Hybrid Plastics, is the most often used monomer to synthesize POSS containing materials, according to the literature.1 It is very well soluble in organic solvents, much better than other commercially available derivatives with Et, c-C5H9, c-C6H11, or Ph groups.1 The solubility and other physical properties such as melting point, crystallization/melting temperature, and thermal stability of R7POSS-Y-MA depend not only on the R groups but also on the Y spacer, connecting the POSS cube with a methacrylate group. Extension of a C3 spacer with a flexible OSiMe2 bridge (OSiMe3C3) provides a compound with a significantly improved solubility and reduced melting point (from 383 to 369 K) and temperature of crystallization (from 378 to 341 K) (Chart 1). However, the thermal stability of OSiMe3C3 was similar to that of C3 monomer (T5% = 502 K, T10% = 515 K). The differences between OSiMe2C3 and C3 compounds can be explained by the lower energy of the crystal lattice of OSiMe2C3, originating from the higher flexibility of the spacer and lower symmetry of the molecule. Solubility of OSiMe2C3 in toluene is so high that stable solutions with 85 wt % (34.8

and can differ substantially in systems with a high dielectric strengths. Dynamic Mechanical Analyses (DMA). DMA were performed using an Advanced Rheometric Expansion System (ARES) equipped with a force-rebalanced transducer. Plate−plate geometry with plate diameters of 13 mm was used. The gap between plates was around 1 mm. Experiments were performed under dry nitrogen atmosphere. Oscillatory shear deformation was applied under conditions of controlled deformation amplitude that was kept in the range of the linear viscoelastic response of the studied samples. Frequency dependencies of the storage G′ and the loss G″ parts of the dynamic complex shear modulus were determined from frequency sweeps measured within the frequency range 10−2−102 rad/s at various temperatures. Master curves for G′ and G″ at a reference temperature were constructed using the time−temperature superposition principle, as discussed in the text. Independently, the isochronal temperature dependencies of G′ and G″ were determined for ω = 10 rad/s and a sweep rate of 2 K/min. Density Functional Theory (DFT). Calculations of the dipole moment of methyl methacrylate (MMA), (i-Bu)7POSS-OSiMe3, (iBu)7POSS(CH2)3-MA (C3), (i-Bu)7POSS-OSiMe2(CH2)3-MA (OSiMe2C3), and (i-Bu)7POSS-OSiMe2(CH2)11-MA (OSiMe2C11) were performed at the DFT-B3LYP16 level of theory and the 6-31G(d,p) basis set. This basis set was employed as a compromise between computational cost and accuracy since computational calculations for its performance on selected MA esters showed a maximum deviation C


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Macromolecules Scheme 1. Polymerization of (i-Bu)7POSS-Y-MA Monomers by ATRP

8.2 × 104 (Mw/Mn = 1.45) was formed. The MW determined for P6 by MALLS was Mn,MALLS = 3.13 × 105 (Mw/Mn = 1.30). Polymers P1−P6 were filtrated through alumina to remove the catalyst and washed with acetone to remove the residual unreacted monomer. Thermodynamics and Local Packing. Figure 1 illustrates the DSC traces of compounds P1, P2, P3, P4, P5, and P6,

mol %) of monomer can be prepared. Thus, 10 g of hybrid silsesquioxane OSiMe2C3 can be dissolved in 1.96 mL of toluene. The milky solution at room temperature (Chart 2, d) becomes transparent when heated to 323 K. Good solubility of POSS-MA is crucial for polymerization, since synthesis of polymers with high polymerization degree by ATRP can only be reached at high monomer concentrations.8,11 Another monomer was synthesized with longer alkyl chain spacer (OSiMe2C11). It has higher thermal stability than the monomers with either C3 or OSiMe2C3 spacers. The 5% and 10% weight loss for OSiMe2C11 were observed at 537 and 611 K, respectively. These temperatures are 308 and 369 K higher than observed for two other monomers. Melting point and temperature of crystallization of OSiMe2C11 are slightly lower than determined for OSiMe2C3. Solubility of OSiMe2C11 in toluene was lower than observed for OSiMe2C3 (Chart 2, e). However, the presence of the OSiMe2 bridge makes that compound better soluble than monomer C3. A relatively simple modification of one of eight groups attached to the POSS cage significantly affected its physical properties. Thus, it is of fundamental interest to explore the effect of spacer length (with C3, OSiMe2C3, and SiMe2C11 spacers) on polymer properties. Polymerization of C3 was carried out with the initial molar ratio 900/1/3/3 of the monomer, EBiB as initiator, CuBr, and PMDETA, respectively. The polymerization of the mixture containing 80 wt % (27.2 mol %) of C3 resulted in the polymer with Mn,GPC = 1.5 × 105 (Mw/Mn = 1.20) (P1) at 85% conversion, within 1 h. The process could not be conducted further due to an ineffective stirring caused by the high solution viscosity. The MW of poly(POSS-MA)s determined by GPC vs linear PMMA standards are strongly underestimated.8,11,20 Therefore, the absolute MW of P1 was measured using multiangle laser light scattering (MALLS) technique. For P1 Mn,MALLS = 4.35 × 105 was almost 3 times higher than Mn,GPC and double compared to the theoretical value (Mn,theor). Then, a series of polymers were prepared by the polymerization of OSiMe2C3 as previously reported.10 Polymers with Mn,MALLS between 1.36 × 104 and 3.257 × 106 (Mn,GPC = 4.8 × 104−5.1 × 105) were obtained. Higher degrees of polymerization than for C3 were obtained due to a better solubility in toluene and higher flexibility of the extended spacer. Polymers P2−P4 were synthesized by batch polymerization while the polymer P5 with the highest degree of polymerization was prepared by semibatch polymerization.8 The batch polymerization of OSiMe2C11 was carried out with the same initial molar ratio of reagents as polymerization P1 and P4. After 1 h at 42% of monomer conversion, the polymer P6 with Mn,GPC =

Figure 1. DSC traces of compounds P1, P2, P3, P4, P5, and P6 obtained on cooling (left) and subsequent heating (right) with 10 K/ min. The second cooling/heating cycle is shown. Notice that the trace of P1, that is lacking the flexible OSiMe2 linker, is featureless.

obtained from the second cooling/heating cycle. With the exception of compound P1 that is lacking the flexible OSiMe2 linker, they show a crystallization/melting process on cooling/ heating with temperatures and heats that depend strongly on spacer length. Clearly, the flexible linker OSiMe2 has an instrumental role in phase formation of polymers bearing POSS moieties. The flexible O−Si bond decouples POSS units from the backbone and facilitates their crystallization. Compounds P2, P3, P4, and P5 melt within the range 353−360 K with a heat of fusion of about 3.3 J/g. In contrast to this, compound P6 with the longer spacer melts at 380 K with a heat of fusion of 11.7 J/g. Both the increased melting temperature and the higher heat of fusion suggest a more ordered structure of P6. To gain more insight into the actual crystal structure, WAXS was performed on the polymers and the corresponding monomers. XRD patterns (Figure 2) of the POSS containing monomers show clear Bragg peaks. The diffractogram obtained for the monomer of compound P6 (i.e., OSiMe2C11) shows qualitative agreement with the patterns obtained for (cyclohexyl)8-POSS, D


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Figure 2. WAXS spectra of compounds P1, P4, P6 (solid blue lines) and of the corresponding monomers C3, OSiMe2C3, OSiMe2C11 (solid red lines) at 293 K. Figure 3. WAXS spectra of compound P4 obtained at different temperatures as indicated. Blue arrows indicate reflections corresponding to POSS crystal unit cell. The red arrow at q = 2.07 nm−1 corresponds to backbone−backbone correlations that improve with increasing temperature.

21

(cyclohexyl)7-POSS-OSiH(CH3)2, and (cyclopentyl)7-POSS(CH2)3-MA.22 However, the peak positions indicate that the lattice constants of P6 are reduced by 16% compared to (cyclohexyl)8-POSS with rhombohedral unit cell (a = 1.157 nm and α = 95.5°).21 This is explained by the size difference between the cyclohexyl and isobutyl groups that protrude from the POSS core with a diameter of 0.56 nm. For the polymers P1 and P4 diffraction peaks are much broader than the corresponding monomers. Amorphous polymers are known to have several broad peaks.23 However, the lack of sharp diffraction peaks makes the assignment of a Bragg spacing (d ∼ 2π/q) from the position of the scattering maximum a crude approximation. Therefore, in amorphous polymers we refer to equivalent Bragg spacing. Poly(n-alkyl methacrylates) sets a good example as they display three such peaks, with the exception of PMMA.24 The first one is the usual van der Waals (VDW) peak due to contacts of atoms. The second, at lower wavevectors, is the low van der Waals (LVDW) peak reflecting interchain correlations. This peak assignment in poly(n-alkyl methacrylates) related to backboneto-backbone correlations is evidenced by plotting the corresponding equivalent Bragg spacing as a function of n1/2, where n is the number of carbon atoms on the alkyl chain.24 This relation can be parametrized as d = d0 + sn1/2, with d0 = 0.7 nm and s = 0.385 nm/CH2. This finding clearly shows an increase in interchain distances with increasing length of alkyl chains. The third peak is an intermediate feature that is close to the single peak in PMMA. These characteristic structural features of poly(n-alkyl methacrylates) confirm the tendency for local segregation as recently reported.24,25 The presence of POSS units connected via spacers to the polymethacrylate backbone further enhances the tendency for local segregation. In contrast to P1 and P5, the widths of the Bragg peaks observed for P6 are comparable to the width found for the corresponding monomer. Comparison of the diffraction patterns indicates that the long and flexible SiMe2C11 spacers in P6 allow POSS units to crystallize similar to the monomers. This observation is in agreement with the DSC measurements (Figure 1) indicating a higher degree of crystalline order in P6 compared to the other polymers. More information on the local packing comes from a temperature-dependent study. Figure 3 shows the diffraction patterns of compound P4 obtained as a function of temperature. Up to a temperature of about 343 K the patterns are dominated by broad peaks at wavevectors 5.77 and 12.3

nm−1 with equivalent Bragg spacing of 1.09 and 0.51 nm, respectively. Some additional superimposed narrower peaks exist at wave vectors 5.77, 7.6, 13.3, and 17.4 nm−1 reflecting the underlying POSS organization. In addition, a weak and broad feature is apparent at a wavevector of 2.07 nm−1. This feature is completely absent in the diffraction patterns of the monomer OSiMe2C3. It originates from backbone-to-backbone correlations having an equivalent Bragg spacing of ∼3 nm. Such distance is consistent with the bulky POSS unit (cage diameter ∼0.56 nm) and spacer groups. The peak intensity increases above the melting temperature of the POSS crystals at ∼353 K. This suggests an interplay of antagonistic interactions favoring POSS crystallization and polymer backbone organization at lower and higher temperatures, respectively. At lower temperatures, POSS units selfassemble and crystallize in a rhombohedral unit cell. This perturbs interchain packing. At higher temperatures, the POSS crystal structure melts. This facilitates backbone-to-backbone correlations that are further improving with increasing temperature. These structural changes are expected to alter the dipole dynamics of pendant POSS cages and of the polymer backbone. Local Dynamics. A comparison of the thermodynamic phase transitions, as seen in DCS, with measurements of the dielectric permittivity, by DS, is shown in Figure 4. The figure compares the DSC traces of compounds P3, P4, and P6 with the corresponding dielectric permittivity and its absolute derivative with respect to temperature, all obtained on heating at a frequency of 104 Hz. It is well-known that ε′(T) provides a sensitive probe of phase transformations for a range of soft materials.26 For better comparison, DSC and DS were carried out with the same heating rate (1 K/min). As expected, both techniques are sensitive to the melting of the rhombohedral unit cell of POSS units (at ∼380 and ∼360 K for compounds P6 and P4, P3, respectively). In addition, a step in dielectric permittivity with a concomitant peak in the derivative representation is evident at ∼293 K for all compounds. An additional minor step is also evident at ∼302 K. These steplike changes in dielectric permittivity reflect the dynamic processes (α and α′) to be discussed below. E


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crystallization temperature. For compound P4, crystallization sets in at 333 K and results in the loss of dielectric strength and broadening of the α′ process. Some representative fits with a sum of two HN functions are shown at T = 323.15 K for both P6 and P4. The parameters characterizing the α′ process are discussed with respect to data shown in Figure 6. The figure depicts the

Figure 4. Top: heat flow obtained on heating with a rate of 1 K/min. Middle: dielectric permittivity obtained on heating at a frequency of 104 Hz. Bottom: absolute derivative of the dielectric permittivity with respect to temperature for compounds P3, P4, and P6. The dashed and dash-dotted lines indicate the respective characteristic temperatures.

Some representative loss spectra of compounds P6 and P4 at several temperatures obtained on cooling are shown in Figure 5. For compound P6, the dielectric loss curves at temperatures above 348 K indicate a single loss mechanism (referred to as α′). At ∼348 K, the α′ mechanism losses intensity and broadens significantly especially at the low-frequency side. In addition, the ionic conductivity shows a discontinuous change at the transition. A second, broader mechanism (called slow DS process) is necessary to describe the loss spectra below the

Figure 6. Temperature dependence of the HN parameters, TΔε (top), m (middle), and mn (bottom), corresponding to the α′-process for compounds P2 (magenta), P3 (blue), P4 (red), P5 (green), and P6 (black). The black and red vertical lines indicate temperatures of POSS crystallization for compounds P6 and P4, respectively.

temperature dependence of the HN parameters, TΔε, m, and mn corresponding to the α′-process for the different compounds. Evidently, POSS crystallization alters significantly the dielectric strength and distribution of relaxation times. Roughly two-thirds of the dielectric strength for the α′ process is lost during POSS crystallization. At the same temperature the process broadens both from the low- (m) and high-frequency sides (mn). At present it is unclear as to which dipoles (MA ester and/or POSS-OSiMe2(CH2)x) contribute to the α′ process. The results from DFT calculations with respect to the monomer dipole moments can elucidate this point (Figure 7). The values of dipole moments indicate that the ester group of MMA dominates the overall contributions to the dipole vectors in all three monomers C3, OSiMe2C3, and OSiMe2C11 since, to a first approximation, these seem to be the result of the vectorial sum between the MMA (μ ∼ 1.65 D) and the (i-Bu)7POSSOSiMe3 (μ = 0.59 D) contributions. Additional contributions, however, are expected from induced moments along the chain and the double bond. Moreover, the vectors adopt almost the same orientation with respect to the oxygen atoms of the MA. The origin of the different dynamic processes can be better discussed with respect to their temperature dependence and the results of DFT calculations. The relaxation times of the different processes in compounds P2−P6 are summarized in Figure 8 using the usual Arrhenius representation. The figure depicts three dielectrically active processes (γ, α′, “slow”), one

Figure 5. Dielectric loss spectra of compound P6 (top) and compound P4 (bottom) at several temperatures within the range from 273.15 to 413.15 K in 5 K steps. Representative fits with a sum of HN functions are shown at T = 323.15 K (the loss curve with the thicker line). The red solid line represents the α′ process and the red dash-dotted line the slower DS process. F


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maximum loss conform to the Vogel−Fulcher−Tammann (VFT) equation:

⎛ B ⎞ τ = τ0 exp⎜ ⎟ ⎝ T − T0 ⎠

(6)

−12

where τ0 (5.0 × 10 s) is the relaxation time in the limit of very high temperatures, B (= 2900 ± 600 K) is the activation parameter, and T0 (= 150 ± 19 K) is the “ideal” glass temperature for this process. The non-Arrhenius temperature dependence of relaxation times, the broadening of the relaxation function, and, most importantly, the sharp loss of dielectric strength at the POSS crystallization temperature (Figure 6) suggest that it reflects the cooperative motion of POSS-OSiMe2 units and of the ester dipoles. This assignment of α′-process is based on the results from DFT calculations revealing that the stronger dipole associate with the MA unit. From the freezing of its dynamics a “POSS glass temperature” can be estimated at ∼247 K. Additional information on the dipole−dipole orientation correlations can be extracted by studying the dielectric strength of α′-process, Δε, at T > 353 K. The static dielectric constant of polar liquids with short-range interactions between molecules has been the subject of the Kirkwood/Fröhlich theory.12a,c The theory has considered an infinite continuum of dielectric permittivity, εS′, and within this a spherical region containing N0 elementary dipoles that were treated explicitly. Based on these assumptions, the dielectric permittivity can be expressed as

Figure 7. Calculated dipole moment (μ) of MMA, (i-Bu)7POSSOSiMe 3 , (i-Bu) 7 POSS(CH 2 ) 3 -MA (C 3 ), (i-Bu) 7 POSS-OSiMe2(CH2)3-MA (OSiMe2C3) and (i-Bu)7POSS-OSiMe2(CH2)11-MA (OSiMe2C11) accompanied by pictorial representation of their optimized structures at the B3LYP level of theory. Basis set employed: 6-31G(d,p). Black arrows represent the dipole moment vectors. Color code of the atoms: oxygen is represented in red color, carbon in cyan, and silicon in yellow. Hydrogen atoms are not shown.

Δε = ε′S − ε∞ =

μ2 N0 1 Fg 3ε0 kBT V

(7)

Here, F = ε′S(ε∞ + 2) /[2(ε′S + ε∞)] is the local field, N0/V is the number density of dipoles expressed as (ρ/M)NA, where ρ is the mass density (obtained assuming the same rhombohedral unit cell as in the monomers resulting in ρ ∼ 1.12 and 1.25 g/ cm3 for OSiMe2C3 and OSiMe2C11, respectively), M is the molar mass, μ is the dipole moment, and g is the Kirkwood− Fröhlich dipole orientation correlation function, g = μ2/μgas2, defined as the ratio of the mean-squared dipole moment measured in a dense system divided by the same quantity obtained in a noninteracting case (gas phase). Employing the DFT results for the gas-phase dipole moments of OSiMe2C3 and OSiMe2C11, we extract g ∼ 0.2. This value is suggestive of a destructive interference of dipoles in the melt state of poly(POSS-MA)s. Dipole orientation correlation values below unity are not uncommon in stereoregular PMMAs (the following sequence is reported g(atactic) < g(syndiotactic) < g(isotactic) < 1).27 Nevertheless this process (α′) is decoupled from the backbone dynamics. The complete freezing of backbone dynamics at the liquid-to-glass temperature is obtained from TMDSC (at T ∼ 280 K) for compounds P2, P3, P4, and P5 and reflects the usual α-process. The latter process is not evident in DS probably for intensity reasons (as most sidegroup dipoles freeze through the faster α′ process). An increase in glass temperature has been reported earlier when POSS units were chemically incorporated directly into the polymer chain. For example, in norbonyl-POSS,5 styryl-POSS,7 and poly(vinylpyrolidone)-POSS hybrid copolymer9 experiments and, in some cases simulation,6 revealed an increased Tg. However, a Tg reduction was reported in poly(acetoxystyryne)-POSS copolymers.9 Interestingly, the liquid-to-glass temperature is 2

Figure 8. Arrhenius relaxation map of the local and global dynamic processes for compounds P2 (magenta), P3 (blue), P4 (red), P5 (green), and P6 (black). The different processes are as follows: (rhombi) γ-process from DS; (squares) α′ process from DS; (up triangles) α process from TMDSC; (down triangles) slow DS process; (filled spheres) chain relaxation from DMA. The black and red vertical lines indicate temperatures of POSS crystallization for compounds P6 and P4, respectively.

process obtained from TMDSC (α), and the chain relaxation obtained from DMA. The γ-process in the glassy state (m = 0.4, mn = 0.2) has an Arrhenius temperature dependence according to ⎛ E ⎞ ⎟ τ = τ0 exp⎜ ⎝ RT ⎠

(5)

−12

with τ0 = 5.6 × 10 s and an activation energy, E, of 36.5 kJ/ mol. This process being insensitive to POSS crystallization is assigned to local rattling motion of POSS cages. The slower and most intense dielectrically process (α′) has a different origin. This process is slightly faster in compound P6with the longer spacerand has a small curvature in the Arrhenius τ(T) representation. The relaxation times at G


Article

Macromolecules ca. 110 K lower than for linear high molecular weight PMMA. The shift in the freezing of the backbone dynamics reflects the combined role of POSS units and flexible spacer. POSS moieties with an open cubic-octameric framework can impart free volume to the system. But more important is the longer and flexible spacer that decouples the POSS cages from the backbone and speed up the segmental dynamics relative to PMMA. The freezing of backbone dynamics, however, takes place at low temperatures, where POSS units are already crystallized to the rhombohedral unit cell. Thus, backbone mobility is already restricted at temperatures where TMDSC identifies the structural (α) process. The higher backbone restriction in the more ordered compound P6 with the longer spacer explains the absence of this feature from the TMDSC curves of P6. Chain Dynamics. Further evidence for the distinctly different local and global dynamics of the poly(POSS-MA)s as compared to linear PMMA chains is provided by dynamic mechanical measurements. Figure 9 gives the temperature

Figure 10. Master curve for the storage modulus (G′, squares) and the loss modulus (G″, circles) for compound P5 using data in the T-range 343 < T < 393 K. The reference temperature was at 343 K. Lines with slopes of 1 and 2 are shown. The master curve for a linear PMMA (Mw = 1.91 × 105 g/mol) is also shown with the open symbols for comparison. The latter curve contains data in the T-range from 403 < T < 473 K, and the reference temperature was at 448 K.

curve at reference temperature of 343 K. At lower temperatures (higher frequencies), at the onset of POSS crystallization, the adhesive contact of the sample with rheometer plates was lost. This explains the absence of the results in this range. At the remaining temperatures, the figure shows a near parallel behavior of the shear moduli with G′ ∼ G″ and a crossing at lower frequencies/higher temperatures to a near terminal behavior. This rheological behavior for compound P5 bearing ∼3300 repeat units (Table 1) is distinctly different from a linear PMMA with ∼1900 repeat units. The latter, despite the smaller number of repeat units, exhibits a rubbery plateau with a modulus of G0N ∼ 0.4 MPa comparable to literature values.28 The absence of a well-pronounced rubbery plateau and the lower values of the shear moduli for compound P5 is attributed to the architecture of polymers bearing POSS units. Evidently, the large fraction of side groups, each bearing a POSS cage unit makes the material not only more mobile but also extremely soft. This situation is reminiscent of supersoft elastomers previously investigated by rheology.29 There, the presence of a large number of dangling chains enhances the material mobility and elasticity, at the same time. In the present case, POSS units impart viscoelasticity in the absence of a low modulus elastic plateau. Future studies will explore the effect of the R group substitution to the POSS units.

Figure 9. Isochronal measurements (ω = 10 rad/s) of the storage (G′; black squares) and loss (G″; red circles) moduli of compound P5 obtained on cooling and subsequently heating with a rate of 2 K/min. Vertical dashed lines indicate the transition temperatures on cooling (blue line) and heating (red line).

dependence of the storage and loss moduli of compound P5 under isochronal conditions (ω = 10 rad/s) on cooling and subsequent heating. The figure depicts the phase transformation from a low temperature elastic phase composed from crystalline POSS units embedded in a “disordered” polymer matrix to a low modulus viscoelastic phase at higher temperatures. More informative on the viscoelasticity of phases for the compound P5 is the frequency dependence of the moduli shown in Figure 10. According to the principle of time−temperature superposition (tTs), the frequency (ω) dependence of the complex shear modulus G* at any temperature can be obtained from a master curve at a reference temperature (Tr) according to the equation G*(ω , T ) = G*(aT ω; Tr)

4. CONCLUSIONS Polymethacrylates with polyhedral oligomeric silsesquioxane (POSS) moieties (poly(POSS-MA)s) containing flexible spacers between the POSS cages and the methacrylate group possess distinctly different properties from their linear counterparts. The pendant POSS moieties modify the existing interchain correlations and further introduce new correlations between the cages. This has drastic effects on the thermodynamic properties as well as on the local and global nanocomposite dynamics. The multiple dynamic processes reflect the relaxation of the side group that extend from the ester dipole to the pendant POSS unit as well as of the polymer backbone. Both processes are cooperative and display a VFT temperature dependence. The freezing of backbone dynamics at the glass temperature is shifted to lower temperatures as

(8)

When the tTs is valid, then at each temperature a single frequency-scale shift factor aT allows superposition of all viscoelastic data at temperature T. The figure depicts the frequency dependence of the real (G′) and imaginary (G″) parts of the shear modulus as measured at several temperatures above the melting temperature of P5 and arranged in a master H


Article

Macromolecules Table 1. Polymerization Conditions and Molecular Weights of Polymers P1−P6 (i-Bu7)POSS-Y-MA polymer

reagents ratioa

mol %

(i-Bu)7POSS-(CH2)3-MA (C3) P1 900/1/3/0/3 27.2 (i-Bu)7POSS-OSiMe2(CH2)3-MA (OSiMe2C3) P2 100/1/2.66/0.66/3.32 14.6 P3 200/1/5.32/1.32/6.64 14.6 P4 900/1/3/0/3 27.2 P5b 3333/1/3/0/3 24 (i-Bu)7POSS-OSiMe2(CH2)11-MA (OSiMe2C11) P6 900/1/3/0/3 27.2

wt %

conv (%)

Mn,theor (×103)

Mn,GPC (×103)

Mn,MALLS (×103)

Mw/Mn,GPC(MALLS)

79

85

721.8

150.0

435

1.20 (1.17)

65 65 80 78

92 83 92 91

93.6 168.9 842.7 3086

48.0 79.0 195.0 510.0

136 242 660 3257

82

44

569.5

80.2

313

a

1.28 1.19 1.40 2.38

(1.07) (1.03) (1.20) (1.24)

1.45 (1.30)

b

[(i-Bu7)POSS-Y-MA]0/[EBiB]0/[CuBr]0/[CuBr2]0/[PMDETA]0. Semibatch polymerization: [OSiMe2C3]0/[EBiB]0 ratio increased from 2000/1 to 3333/1. C.; Xin, Y.; Wang, H.-F.; Shi, A.-C.; Newkome, G. R.; Ho, R.-M.; Chen, E.-Q.; Zhang, W.-B.; Cheng, S. Z. D. Proc. Natl. Acad. Sci. U. S. A. 2013, 110, 10078−10083. (c) Yue, K.; Liu, C.; Guo, K.; Wu, K.; Dong, X.-H.; Liu, H.; Huang, M.; Wesdemiotis, C.; Cheng, S. Z. D.; Zhang, W.-B. Polym. Chem. 2013, 4, 1056−1067. (3) (a) Mather, P. T.; Jeon, H. G.; Romo-Uribe, A.; Haddad, T. S.; Lichtenhan, J. D. Macromolecules 1999, 32, 1194−1203. (b) Bharadwaj, R. K.; Berry, R. J.; Farmer, B. L. Polymer 2000, 41, 7209−7221. (4) Haddad, T. S.; Lichtenhan, J. D. Macromolecules 1996, 29, 7302− 7304. (5) Lee, A.; Lichtenhan, J. D. Macromolecules 1998, 31, 4970−4974. (6) Xu, H.; Kuo, S.-W.; Lee, J.-S.; Chang, F.-C. Macromolecules 2002, 35, 8788−8793. (7) Li, G. Z.; Wang, L.; Toghiani, H.; Daulton, T. L.; Koyama, K.; Pittman, C. U. Macromolecules 2001, 34, 8686−8693. (8) Franczyk, A.; He, H.; Burdyńska, J.; Hui, C. M.; Matyjaszewski, K.; Marciniec, B. ACS Macro Lett. 2014, 3, 799−802. (9) (a) Wang, J.-S.; Matyjaszewski, K. J. Am. Chem. Soc. 1995, 117, 5614−5615. (b) Kato, M.; Kamigaito, M.; Sawamoto, M.; Higashimura, T. Macromolecules 1995, 28, 1721−1723. (c) Matyjaszewski, K.; Xia, J. Chem. Rev. 2001, 101, 2921−2990. (d) Gao, H.; Matyjaszewski, K. Prog. Polym. Sci. 2009, 34, 317−350. (e) Matyjaszewski, K.; Tsarevsky, N. V. Nat. Chem. 2009, 1, 276−288. (f) Matyjaszewski, K. Macromolecules 2012, 45, 4015−4039. (g) Hui, C. M.; Pietrasik, J.; Schmitt, M.; Mahoney, C.; Choi, J.; Bockstaller, M. R.; Matyjaszewski, K. Chem. Mater. 2013, 26, 745−762. (h) Matyjaszewski, K.; Tsarevsky, N. V. J. Am. Chem. Soc. 2014, 136, 6513− 6533. (10) Pyun, J.; Matyjaszewski, K. Macromolecules 1999, 33, 217−220. (11) Raus, V.; Č adová, E.; Starovoytova, L.; Janata, M. Macromolecules 2014, 47, 7311−7320. (12) (a) Kremer, F.; Schoenhals, A. Broadband Dielectric Spectroscopy; Springer: Berlin, 2002. (b) Floudas, G.; Paluch, M.; Grzybowski, A.; Ngai, K. L. In Molecular Dynamics of Glass-Forming Systems; Springer: Berlin, 2011. (c) Floudas, G. In Polymer Science: A Comprehensive Reference; Matyjaszewski, K., Mö ller, M., Eds.; Elsevier BV: Amsterdam, 2012; Vol. 2.32. (13) Havriliak, S.; Negami, S. Polymer 1967, 8, 161−210. (14) Williams, G.; Watts, D. C. NMR: Basic Princ. Prog. 1971, 4. (15) Jäckle, J.; Richert, R. Phys. Rev. E 2008, 77, 031201. (16) (a) Lee, C.; Yang, W.; Parr, R. G. Phys. Rev. B 1988, 37, 785− 789. (b) Becke, A. D. J. Chem. Phys. 1993, 98, 5648−5652. (17) Hickey, A. L.; Rowley, C. N. J. Phys. Chem. A 2014, 118, 3678− 3687. (18) Frisch, M. J., et al. Gaussian 03, Revision E.01; Gaussian, Inc.: Wallingford, CT, 2004. (19) Humphrey, W.; Dalke, A.; Schulten, K. J. Mol. Graphics 1996, 14, 33−38. (20) Hirai, T.; Leolukman, M.; Jin, S.; Goseki, R.; Ishida, Y.; Kakimoto, M.-A.; Hayakawa, T.; Ree, M.; Gopalan, P. Macromolecules 2009, 42, 8835−8843.

compared to linear PMMA. At the same time, the nanocomposites appear softer than linear PMMA chains of similar degrees of polymerization under isofriction conditions. Results presented here suggest that POSS cages can be readily employed as nanometer size blocks that when coupled to a polymer backbone via flexible spacers impart free volume, mobility, and control over the mechanical behavior of nanocomposites.

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ASSOCIATED CONTENT

S Supporting Information *

Experimental details. The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.5b00663.

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AUTHOR INFORMATION

Corresponding Author

*E-mail gfl[email protected] (G.F.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank Andreas Hanewald and Christine Rosenauer for the technical assistance. This work was cofinanced by the E.U.-European Social Fund and the Greek Ministry of Development-GSRT in the framework of the program THALIS. The current work was supported by the Research unit on Dynamics and Thermodynamics of the UoI cofinanced by the European Union and the Greek state under NSRF 20072013 (Region of Epirus, call 18). Support by the European Regional Development Fund: Project No. UDAPOIG. 01.03.01-30-173/09, VENTURES Programme, funded by The Foundation for Polish Science (FNP) is acknowledged. The support by the National Science Foundation (DMR 1501324) is also acknowledged. The authors gratefully acknowledge the computing time granted by the Research Center of Scientific Simulations, University of Ioannina.

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REFERENCES

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Polymethacrylates with Polyhedral Oligomeric Silsesquioxane (POSS

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