Reduced Phase Separation and Slowing of Dynamics in

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Reduced Phase Separation and Slowing of Dynamics in Polyurethanes with Three-Dimensional POSS-Based Cross-Linking Moieties Konstantinos N. Raftopoulos,*,† Stefanos Koutsoumpis,‡ Małgorzata Jancia,† James P. Lewicki,§ Konstantinos Kyriakos,∥ Harris E. Mason,§ Stephen J. Harley,§ Edyta Hebda,† Christine M. Papadakis,∥ Krzysztof Pielichowski,† and Polycarpos Pissis‡ †

Department of Chemistry and Technology of Polymers, Cracow University of Technology, ul. Warszawska 24, 31-155 Kraków, Poland ‡ Department of Physics, National Technical University of Athens, Iroon Polytechneiou 9, Zografou Campus, 157 80, Athens, Greece § Lawrence Livermore National Laboratory, 7000 East Ave., Livermore, California 94550, United States ∥ Physik-Department, Fachgebiet Physik weicher Materie, Technische Universität München, James-Franck-Str. 1, 85748 Garching, Germany S Supporting Information *

ABSTRACT: Octa-OH-functional POSS has been incorporated into a model polyurethane elastomer as a comparatively massive and notionally “robust” 3-dimensional cross-linking core. The effects of this cross-linking moiety on the morphology and molecular dynamics of the system are studied over a range of size and time scales. Microscopy, scattering, spectroscopic, thermal, and dielectric techniques, in agreement with each other, show that the covalent inclusion of the crosslinking particles restricts microphase separation, inhibits the formation of hard-block domains, and decelerates the motional dynamics of the polyurethane backbone. The effects on both the morphology and the dynamics of the polyurethane system are not continuous but occur in a steplike manner in the loading region of 4−6 wt % POSS. This critical region is thought to correspond to a sterically induced transition from one dominant morphology (microphase segregated) to an increasingly homogeneous nanophase segregated domain morphology. Contrary to expectations, cross-linking, even by the presumably rigid siliceous nanoparticles, reduces the mechanical modulus. In conjunction with the reduction of microphase separation, this observation indicates that the hard microdomains reinforce the polymer more effectively than the chemical cross-links.

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INTRODUCTION The so-called nanobuilding block approach facilitates dispersion and targeted placement of nanoparticles in polymers. By this approach, properly functionalized nanoparticles react with functional groups on the macromolecular chain or with monomers or oligomers during synthesis and are effectively bound on the macromolecular structure itself. Polyhedral oligomeric silsesquioxanes (POSS) are ideal nanobuilding blocks due to the versatility in the reactive groups they may possess. POSS typically consists of a siliceous polyhedral core with Si vertices and Si−O−Si edges surrounded by a shell of organic chains attached to the vertices. The size and complexity of these organic vertex groups vary from a single H atom to oligomers with Mw of a few hundred g/mol. Virtually any functional group may be present on these groups, including alcohol, amine, ether, isocyanate, etc., as comprehensively described in numerous review articles.1−6 The number of reactive vertex groups determines to a large extent the topology of the host polymer chain. Moieties without © XXXX American Chemical Society

reactive groups are typically blended in the polymer matrix, much like in conventional nanocomposites, but the organic ligand facilitates solubility and dispersion.7,8 Particles with one reactive vertex group are typically pendent on9 or end-cap10 the chain. Particles with two reactive groups lie along the chain contour in a beadlike configuration.11 Finally, particles with more reactive vertex groups act as comparatively massive, threedimensional chemical cross-links. Several studies following different chemical approaches report on polyurethane and polyurea systems with POSS acting as chemical cross-links: octaaminophenyl-POSS crosslinking a typical polyurethane copolymer,12 octavinyl-POSS cross-linking an acryl end-capped polyurethane prepolymer,13 trisilanolphenyl-POSS cross-linking a polyurea prepolymer,14 octaisocyanate-POSS cross-linking a polyurethane network.15 Received: November 15, 2014 Revised: February 6, 2015

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DOI: 10.1021/ma5023132 Macromolecules XXXX, XXX, XXX−XXX

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Synthesis of PU/Octa-OH-POSS Hybrids. The polyurethane/octaOH-POSS hybrid elastomers were prepared with 0, 2, 4, 6, 8, and 10 wt % octa-OH-POSS. The hybrid materials were synthesized in batches of 30 g, with composition reported in Table 1. A two-step

Increases in Tg of POSS-modified PU systems have been reported and are attributed to nanoreinforcement and restriction of mobility due to constraints imposed by crosslinking.12,13 In the case of PU−POSS network systems, particle aggregation appears to partially counteract the reinforcing effect of the POSS cross-linker.15 Casalini et al. have highlighted the fact that POSS molecules function as chemically reactive additives, rather than as conventional nanofillers,14 and indeed, the view that POSS is a form of nanosilica is, in general, losing traction in the literature.6,16 Both glassy and rubbery moduli tend to increase with POSS content due to reinforcement and cross-linking,12−14 but above a critical loading, aggregation12,13 and steric disruption of the existing PU domain microstructure by POSS16 counteract the stiffening effect. The underlying mechanisms that control molecular mobility and mechanical properties are complex and not always clear in such multicomponent systems, but one should consider nanoreinforcement, changes in the network topology, and of course effects on microphase separation.6 Which of those mechanisms is prevalent will depend strongly on the chemical nature of the components that constitute each system and the subsequent mutual interactions. In previous studies, we have examined in detail the morphology and dynamics in polyurethane POSS hybrids with particles either pendant16−19 or along the macromolecular chain.20 The matrix in these studies is a simple model polyurethane with poly(tetramethylene ether glycol) (PTMEG) as the soft component and methylene diphenyl diisocyanate (MDI) as the hard component. The chain is extended by butane diol, which is partly substituted by OH functional POSS. In all cases, we observed a moderate deceleration of PU dynamics, but it arose from substantially different mechanisms. In the pendant systems, it was reinforcement by the particles combined with disruption of microphase separation. Interestingly, reinforcement was more prevalent in materials where the soft and hard segments had relatively large Mw, while reduced separation in materials with shorter (low Mw) segments.18 In the case of the beadlike structure, PU extended by POSS mostly phase-separated from the matrix and showed its own segmental dynamics, at a longer time scale than the “unmodified” phase by approximately 3 orders of magnitude. In this current study, we investigate the use of octa-OHfunctional POSS as a comparatively massive and notionally “robust” 3-dimensional cross-linking core in the same model MDI-based phase-segregated PU elastomer and study the changes in global morphology and molecular dynamics of the system as a whole through a combination of microscopy, X-ray scattering, nuclear magnetic resonance, thermal analysis, and dielectric relaxation spectroscopy.

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Table 1. Masses of Reactants Used in the Synthesis of the Hybrid Elastomers reactant mass (g) POSS content (wt %)

MDI

PTMG

octa-OH-POSS

0 2 4 6 8 10

11.80 11.34 10.74 10.14 9.54 8.94

15.00 15.00 15.00 15.00 15.00 15.00

0.00 0.60 1.20 1.80 2.40 3.00

procedure was followed (Figure 1): MDI was charged into a 100 mL three-necked round bottomed reactor equipped with a mechanical stirrer, a thermometer, and an argon inlet, heated to 70 °C with a thermostated electrical heating mantle, and a solution of octa-OHPOSS in 2 mL of THF was added. The reaction was carried out under inert atmosphere at 80 °C for 2 h to form MDI-octa-OH-POSS “prepolymer” and to evaporate the THF. In the second stage, the “prepolymer” was mixed with suitable amounts of PTMG and 1,4butanediol. PTMG accounted for the 50 wt % of the final material, and the amount of butanediol was determined by a functional groups equivalents determination21 to allow for complete reaction of still free MDI units. This process also revealed that the OH groups of POSS reacted successfully with the MDI monomers. The resultant, still uncured, viscous mixture was poured out onto an open steel mold, and cured at 80 °C for 18 h to form a solid elastomer. Atomic Force Microscopy. Atomic force microscopy (AFM) images were recorded on the reference matrix and selected hybrids with a Veeco diInnova microscope in a tapping mode, on surfaces produced by a microtome at room temperature. X-ray Scattering. Small-angle X-ray scattering (SAXS) measurements were performed on a Ganesha 300XL SAXS-WAXS system (SAXSLAB ApS, Copenhagen/Denmark) equipped with a GENIX 3D microfocus X-ray source and optic together with a three-(scatterless)slit collimation system. Both the sample chamber and the beam path were under vacuum. A two-dimensional (2D) Pilatus 300 K detector was used, which can move to the desired sample-to-detector distance (SDD). The Xray source was operated at 50 kV/0.6 mA with a Cu anode (Kα = 1.542 Å). Silver behenate was used as reference for the angular calibration. Three SDDs were chosen (106.2, 406.2, and 1056.2 mm), resulting in a q-range from 0.005 to 2.5 Å−1. The acquisition time varied for the different SDD’s; i.e., it was 900, 1800, and 3600 s for 106.2, 406.2, and 1056.2 mm, respectively. A pin diode was used to measure the transmission of each sample taking as reference the intensity of the empty beam. All images were corrected for cosmic background and parasitic scattering. No further background correction was required since the sample was clamped on a generic sample holder without interference by any mica windows or glass container. The obtained 2D images were azimuthally averaged. Finally, the intensity curves were normalized over the thickness of the samples. Nuclear Magnetic Resonance. All 1H spin diffusion studies were performed on a 300 MHz Techmag Apollo spectrometer at an operating frequency of 301.13 MHz. The samples were contained in 4 mm o.d. rotor and held statically within a 4 mm Bruker HX probe configured for 4 mm o.d. rotors. Standard inversion recovery techniques were employed to measure 1H spin−lattice relaxation constants (T1), and the Goldman−Shen22 pulse sequence was employed in order to measure the rate of magnetization transfer from hard to soft segments, using a 30 μs T2 filter and allowing the 1H magnetization to diffuse during a second delay period. Individual

EXPERIMENTAL SECTION

Materials and Synthesis. Materials. 4,4′-Diphenylmethane diisocyanate (MDI) was purchased from Sigma-Aldrich and used as received. Poly(tetramethylene glycol) (PTMG) (Terathane) with average molecular weight 1400 g/mol was provided by Invista, 1,4butanediol was obtained from Sigma-Aldrich, and octa(3-hydroxy-3methylbutyldimethylsiloxy) POSS (octa-OH-POSS) was purchased from Hybrid Plastics Ltd. Polyether diol, chain extender, and octa-OHPOSS were dried under vacuum at 80 °C for 10 h before the reaction. Anhydrous tetrahydrofuran (THF) was purchased from POCH and used as received. B


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Figure 1. Reaction route and sketch of the anticipated complex network. Some links have been omitted for clarity. Fredholm integral equations. Many algorithms exist to accomplish this transformation, including the fast Tikhonov regularization routines, CONTIN28 and FTIKREG29 that only vary in the method used to find the criteria for the regularization parameter. In this work, all T2 data were processed using the FTIKREG routine. Modulated Differential Scanning Calorimetry. Modulated differential scanning calorimetry (MDSC) applies a harmonic modulation on a typical DSC linear ramp and then by a signal analysis process deconvolutes the typical DSC thermogram in the socalled reversing and nonreversing components.30,31 The reversing component includes phenomena that are reversible in the time scale of modulation such as the glass transition step and melting of outer layers of crystallites. The nonreversing components contain the remaining phenomena, such as melting of crystal cores, enthalpy relaxation accompanying the glass transition, and any chemical reactions. MDSC experiments were performed on a TA Q200 calorimeter, cooled with liquid nitrogen, purged with helium, and calibrated with indium standards. The apparatus is also equipped with Tzero functionality which compensates for resistance and capacitance imbalances and considerably improves baseline.32 Specimens of 6− 10 mg were measured in the range of −160 to 250 °C at rate 5 K/min and a modulation of ±3 K over a period of 60 s. Glass transition temperatures (Tg) and heat capacity steps (ΔcP) were quantified at the midpoint.

spectra were collected as a function of the second delay in 28 steps from 0.1 ms to 10 s for 16 acquisitions each with a 6 s pulse delay. All other T2 relaxometry studies were performed on a low field Bruker Minispec spectrometer at 37 ± 0.1 °C under static conditions. 90° pulse lengths of τP = 2.25 μs and recycle delays of 15 s were used. Highly rigid domains relax more rapidly than the dead time of the probe; therefore to regain information that may be lost during this time, a special form of refocusing sequence known as the magic sandwich echo (MSE)23 was employed. With MSE, near-quantitative refocusing of fast relaxing components is achieved where subsequent application of a Carr−Purcell−Meiboom−Gill (CPMG) train removes the effects of magnetic field inhomogeneities.24 This technique is highly effective at obtaining free induction decays (FIDs) that are representative of large distributions of relaxation times. The application of the MSE yields an FID that is a superposition of decaying exponentials whose amplitudes and time constants are representative of unique molecular motion distributions within the material. However, the extraction of this information is mathematically nontrivial.25,26 Direct nonlinear regression to an unknown sum of exponentials is possible yet impractical as it requires a priori knowledge of the relaxation behavior of the material of study. Ideally, a mathematical transform is sought to deconvolute the time domain data into a spectral domainmuch akin to a fast Fourier transform. Such a spectral function can only be found upon solution of the Laplace integral equation27 which is part of a more general class of C


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Figure 2. AFM phase images at 20 × 20 μm (top) and 500 × 500 nm (bottom) magnification. Arrows point at structures which are described in detail in the text. Dynamic Mechanical Analysis. Dynamic mechanical analysis (DMA) thermograms were recorded with a Netzsch 242C dynamic mechanical analyzer in the temperature range of −110 to 120 °C at 2 K/min. Frequencies of 1−50 Hz were used. Samples were beams of approximate size 22 × 7 × 2 mm3, cut from the original thick film. They were measured in a three-point bending configuration on a 10 mm sample holder. Thermally Stimulated Depolarization Currents. Thermally stimulated depolarization currents (TSDC) is a special dielectric technique in the temperature domain. It roughly corresponds to measuring isochronally dielectric loss at a low equivalent frequency of the order of mHz,33 typically not accessible by conventional dielectric techniques. In TSDC, a film of the material is placed between the plates of a parallel capacitor and stabilized to a polarization temperature TP. Then a dc field E of the order of kV/mm is applied for a polarization time tP. The system is then cooled down to a low temperature, the field is switched off, the capacitor is short-circuited through a sensitive electrometer, and the specimen is heated at a constant rate bheat. As the relaxation times of the molecular mobility mechanisms decrease, the material is depolarized and the resulting current I is recorded by the electrometer. For comparison purposes, I is subsequently normalized as34

Inorm =

I Id = EA VA

temperature was controlled by a Novocontrol Quatro cryosystem. The frequency range was 10−1−106 Hz, and the temperature range was −130 to 200 °C at steps of 5 or 10 K. The results were quantified in terms of complex dielectric function ε*. More details of this wellestablished method may be found in refs 35−37.

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RESULTS AND DISCUSSION Atomic Force Microscopy (AFM). Figure 2 shows AFM phase images on microtomed surfaces of the matrix and selected hybrids. In this type of imaging, stiff areas appear brighter than flexible ones. The upper row of Figure 2 shows images at low magnification. Starting with the unmodified matrix, globular structures of 3−6 μm with a microphase separated interior can be distinguished. These globular structures are dispersed in a smoother, i.e., less phase separated, continuous phase. This is illustrated in more detail in Figure S1a of the Supporting Information, which shows a close-up of the border. Inside the globule, hard domains (bright areas, bottom row of Figure 2) form elongated structures with diameters of 30−40 nm. Proceeding to the 4 wt % hybrid, the globules are observed to shrink and assume a more spherical shape. The hard elongated structures inside the globules show a reduction in thickness (10−20 nm). In addition, spherical hard structures of ca. 45 nm with a soft interior are observed (Figure 2, point B1). We interpret these new structures as small numbers of POSS moieties interconnected with MDI links, forming the soft core surrounded by hard segments. Note that although the siliceous POSS core is considered as a stiff moiety, the surrounding organic chains, occupying most of the particle volume are rather flexible. Larger (100−300 nm) particles are also observed to be dispersed in the soft phase (Figure 2, point B2, and Figure S1b). These particles may also relate to a POSS-rich phase. In the 6 wt % hybrid, the highly phase separated globules are reduced in diameter, yet numerous, and no particles are visible in the continuous phase. Inside the globules, the hard elongated

(1)

where A is the area of the plates, d the sample thickness, and V the polarizing voltage. The specimens in this investigation were films of thickness ca. 1.5 mm, cut from the original thick film, and placed between polished brass electrodes of diameter 20 mm. Experimental parameters were TP = 25 °C, V = 950 V (E ≈ 1 kV/mm), tP = 5 min, bcool = 10 K/min, and bheat = 3 K/min. The experimental setup comprised a Keithley 617 electrometer, a homemade voltage source, and a Novocontrol Quatro cryosystem to control the temperature. Dielectric Relaxation Spectroscopy. In this technique, an alternating voltage, over a broad frequency range, is applied to a capacitor filled with the sample. The complex impedance is then measured by a dielectric response analyzer. All measurements were performed on a Novocontrol Alpha analyzer, using the same specimens as with TSDC experiments. The D


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A broad amorphous halo in the range 0.7−1.8 Å−1 is typical of MDI based44−48 and other49 polyurethanes. Thus, it is characteristic of the hard segments and possibly related to short-range order induced by hydrogen bonds between urethane groups, which drive microphase separation.47 This halo comprises at least two subpeaks, similar to reflections by different types of MDI-butanediol crystals.44 Nevertheless, there are no sharp peaks that would indicate actual hard segment crystallinity. For the low loading materials (2−4 wt %), a weak peak emerges at 0.7 Å−1. Koberstein and Galambos attribute this peak to c-axis reflections of the MDI-butanediol crystals,44 but in our earlier work, it has also been correlated to hard segments dispersed in the soft phase.48 For the high loading materials, a similar peak occurs at slightly lower q, but in this case, a contribution by thepresumablybroadened main POSS reflection should be considered. The small- and wide-angle regions of the X-ray diffractogramroughly separated at ca. 0.1 Å−1correspond to substantially different structures of the material, i.e., microphase separated domains and order inside those domains, respectively. The typical quantification approach differs for these types of phenomena. It is noteworthy that the two regions have never been measured in a single run by the same instrument. Thus, we base our analysis to common practice and perform the quantification in two stages, starting with the fitting of appropriate model functions to the wide angle side (high q). Wide-Angle Region−Short-Range Order. The low q diffractogram is known to diminish according to the Porod− Ruland theory as39,50

domains are thicker as compared to the matrix and the 4 wt % hybrid. Spheroidal structures (Figure 2, points C1 and C2) again possibly related to POSSare visible here, too, but the soft core is more difficult to distinguish. Note the significant difference in morphology between 4 and 6 wt % hybrids. In the following sections, this abrupt change and its reflection to the other properties of the materials will be discussed. The 10 wt % hybrid has a globular structure similar to that of 6 wt % one. In addition, large (ca. 30−200 nm) “sacs” with a soft core surrounded by a stiff shell are present on all of its surface as a result of the high POSS loading (Figure 2, point D1, and Figure S1c,d). Tan et al. observed similar “sacs” in solutions of star polymers with the same POSS moiety as core and attribute them to solvophobic or stereocomplex driven aggregation.38 Here, we may interpret this form similarly to the structures in the 4 wt % hybrid, but now with significantly more POSS moieties participating in the core, in a structure similar to that proposed by Tan et al.38 X-ray Scattering. The X-ray diffractograms of the elastomers consist of several overlapping regions (Figure 3) which will be described in the order of increasing scattering vector q.

lim I(q) =

q →∞

KP exp( −σ 2q2) + Ib q4

(2)

where KP is a constant related to the surface to volume ratio of the phases, Ib is related to the thermal fluctuations, and σ is a measure of the thickness E of the interfacial region (E ≈ √12σ). Thus, a term of the form of eq 2 was used as a background to the high q region (q > 0.1 Å−1). We modeled the main amorphous halo by a sum of three Lorentz peaks. One more Lorentz term was used for the small peak around 0.7 Å−1. The wide halo which we relate to the octa-OH-POSS could not be modeled by a Lorentz term because of its large width; however, a Gauss term provided an adequate fitting. An example of a fitted curve with its respective components is provided in the Supporting Information (Figure S2). Following this approach, we quantified the position qmax and width of the peaks (Figure 4). The positions of the three hard domain (HD) peaks, corresponding to the intersegmental distances of hard segments (4.3, 4.9, and 5.5 Å), are practically unaffected by the inclusion of POSS into the matrix. We assume that at least part of the hard domains are not penetrated by the particles and therefore retain their native structure. Nevertheless, POSS inclusion broadens the respective peaks (results shown in the Supporting Information, Figure S3). This indicates either a broader distribution of the interdomain distances or, more likely, shrinking of the respective structures. The POSS related halo around 0.4 Å−1 moves toward smaller angles (larger length scales, Figure 4). This indicates an increase in the average distance between the particles. The qmax of the “prepeak” around 0.7 Å−1 undergoes a steplike decrease between 4 and 6 wt % (Figure 4). Because of

Figure 3. X-ray diffractograms recorded with the matrix and the hybrids, corrected for transmission and sample thickness. The WAXS diffractogram of octa-OH-POSS is also included for comparison. A small arrow indicates the position of the prepeak, to be discussed in the text.

A prominent peak at q ≈ 0.04 Å−1 is related to the microphase separation and, as such, observed in several polyurethanes.39−42 Interestingly, the peak consists of two componentsan observation that, to the best of our knowledge, is reported only in ref 42. Before we proceed to any analysis procedure, we would like to point out that the intensity of the diffractograms in this region obviously decreases by an order of magnitude. This is a straightforward indication that cross-linking by POSS suppresses the microphase separation. Around q = 0.4 Å −1 , an amorphous halo emerges systematically on the addition of POSS and is clearly related to their presence within the matrix. Interestingly, the position of the halo coincides with a characteristic reflection of the neat crystalline nanoparticles. A similar change of a crystalline peak to a halo has been reported when rather short PEG chains were attached to a highly crystalline octa(dimethylsiloxy)-POSS43 and may be attributed to small or very distorted crystalline domains. E


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Figure 4. Position qmax of the X-ray peaks in the high q region. The spacings d, corresponding to selected q-vectors, are annotated. The error bars are often smaller than the corresponding symbol.

Figure 5. 1-D correlation functions as calculated for all samples under investigation (top) and an appropriate zoom-in (bottom).

Table 2. Morphological Characteristics of the Matrix and the Hybrids As Measured by SAXSa

the expected contribution of the main POSS reflection, the interpretation of this step is nontrivial. Nevertheless, an abrupt change of morphology in the soft phase is justified, in consistency with AFM observations. Finally, with respect to the Porod−Ruland background, it is interesting to point out that the parameter σ, related to the thickness of the interfacial regions, consistently converged to practically 0 (≈10−9 Å) for all materials. This suggests that the boundaries between hard and soft domains are relatively sharp. Small-Angle Region−Long-Range Order. For the quantitative analysis of the small-angle region, we used the “autocorrelation triangle” approach based on the normalized correlation function,48 in detail described in ref 51. By this approach, a model with a pseudo-two-phase structure is assumed: a hard phase consisting of hard domains dispersed in apartially cross-linkedsoft matrix consisting of PTMEG, non-phase-separated MDI-butanediol structures, and the POSS cross-linker. We justify the assignment of POSS to the soft phase, based on the WAXS observation that the intersegmental distances in the hard domains remain practically unaffected upon addition of POSS. The one-dimensional correlation function γ1(r), where r is the length scale in real space, is defined as39,48,51 ∞

γ1(r ) =

SAXS LP1 (nm)

LHD (nm)

LSP (nm)

α

Q/Q(PU)

1/2 (ts,0 m) (s1/2)

drigid (nm)

0 2 4 6 8 10

9.3 10.0 9.6 9.6 9.4 9.1

3.5 3.2 2.5 1.9 1.7 1.8

5.8 6.8 7.1 7.7 7.8 7.4

0.37 0.32 0.26 0.20 0.18 0.19

1 0.94 0.79 0.51 0.44 0.29

7.36 7.10 6.40 4.01 3.61 3.10

22.3 21.5 19.4 12.1 10.9 9.4

LP1 = primary long period, LHD = thickness of hard domains, LSP = thickness of soft phase, α = volume percentage of hard domains, Q/ Q(PU) = relative invariant with respect to the value of the matrix and spin-diffusion derived calculated effective hard-block domain size (drigid).

The LP2 maximum is not clearly discerned, but it may be observed that it moves progressively to lower distances, indicating an approach of the corresponding structures. Finally, we would like to point out that the denominator in eq 3 is the invariant Q=

∫0 I(q)q cos(qr ) dq 2

∫0 I(q)q dq

POSS content (wt %)

a

2

∞

NMR

∫0

∞

I(q)q2 dq

(4)

which is a measure of the degree of microphase separation (DMS). A strict calculation of DMS requires the assumption of a two-phase morphology and the calculation of electron densities of the components, which here is nontrivial and beyond the scope of the current study. Nevertheless, we have calculated the invariants and normalized them with the value for the matrix (Table 2). The invariant decreases monotonically down to 29% of the value of the matrix for the highest loading, confirming a strong suppression of the microphase separation. The most prominent reduction occurs between contents 4 and 6 wt %, indicating an abrupt morphological change in agreement with AFM and the results from the analysis of the WAXS region. Nuclear Magnetic Resonance (NMR). Spin-Diffusion Analysis of the Rigid Phase Domains. Solid-state spindiffusion experiments are an effective means of determining the domain sizes in heterogeneous polymers, including phasesegregated polyurethanes.52 It is especially amenable to the analysis of rigid phase domains below Tg within a more “mobile” matrix.53 Here we have applied the classic Goldman−

(3)

For the calculation of γ1(r) we subtracted from the diffractogram the wide-angle peaks as calculated in the previous paragraph. Extrapolation to high angles followed Porod’s equation with the parameters that were calculated in the previous fit. The calculated γ1(r) is shown in Figure 5. The first maximum around 100 Å corresponds to the dominant component of the SAXS peak and may be referred to as the long period (LP1) of the materials, i.e., the average distance between two neighboring hard domains. A minor maximum (LP2) enhanced on addition of POSS lies around 140 Å. The LP1 decreases on addition of POSS (Table 2). Assuming a pseudo-two-phase morphology and following the autocorrelation triangle approach,51 we calculate the average thickness of the hard (LHD) and soft (LSP) domains as shown in Table 2. Data indicate, much like the AFM results, that upon addition of POSS the hard domains become thinner and occupy less volume in the material. F


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Macromolecules Shen22 pulse sequence to the analysis of the POSS−PU hybrid materials and measure the rate of transfer of magnetization from the fast-relaxing 1H spin population associated with the rigid, hard-block domains to the surrounding soft-block phase. Despite not being a directly quantitative analysis technique with respect to absolute domain size, spin-diffusion experiments such as these can be a convenient method of probing effective regions of varying mobility within a PU system as a whole. Shown in Figure 6 are the spin-diffusion curves obtained for all six systems studied.

in itself strong evidence of a decrease on the size (increase in the surface area) of these phase domains. What is also interesting is that the trend does not scale linearly with POSS mass fraction; rather, an abrupt increase in exchange rates between 4 and 6 wt % POSS is observed. Such a change over a comparatively narrow window is consistent with a shift in the hard segment phase domain from one dominant morphology to anotheras was observed in both the AFM and SAXS data. If we apply a diffusional model to the data53 and assume a spherical geometry for the hard-block domains, we can use the measured magnetization exchange rates to estimate the effective average diameter of the hard segments using the equation drigid =

2ε π

Dtms,0

(5)

where drigid is the effective domain size (in nm), D is the spindiffusion coefficient (0.8 nm2/ms for a rigid system below Tg), and ε is is the number of orthogonal directions relevant for the spin diffusion process where ε = 3 for a system with discrete phases. Based on this model, values of the effective measured domain size as determined by spin-diffusion for each material have been calculated and are tabulated in Table 2. From the results shown in Table 2 it is clear that as the mass fraction of POSS increases, the effective domain size of those protons within the samples considered to be “rigid” with respect to the material as a whole (therefore assigned as the hard-block domains) decreases significantly with increasing mass fraction of POSS. The sizes of the “rigid domains” as measured by NMR are significantly larger than those determined through SAXS measurements; however, the size of the NMR measured rigid phases are in relatively good agreement with the interglobular structures reported in the AFM analysis. It appears that the spin diffusion analysis has therefore distinguished defined regions of reduced mobility throughout the matrix which, although may not be formally fully crystalline in makeup, are nevertheless part of a global segregated domain structure. These rigid domains are observed to decrease by over 50% in diameter between 0 and 10 wt % POSS inclusion levela trend which correlates strongly with both the SAXS and AFM data. This reduction in rigid domain size has been observed by the authors previously in related POSS−PU16 and PU−carborane54 hybrid systems and is thought to be a result of the influence of the sterically bulky POSS cages disrupting the existing microphase separated morphology which is replaced by a new, nanophase separated domain structure. Magic Sandwich Echo (MSE) Analysis of POSS−PU Hybrid Systems. While spin-diffusion techniques are most useful and convenient for the analysis of rigid phases below Tg in polymer systems, mobile interfacial and elastomeric phases present more a challenge to diffusional methods. We have therefore applied T2 relaxometry for the analysis of the elastomeric regions of POSS−PU hybridsgiven that values of the transverse relaxation time, T2, can be broadly related to mobility (T2 ∝ “proton mobility”). As in previous studies,16,55 data from MSE experiments have been submitted to a regularization routine (see Experimental Section) in order to derive distributions of T2 values for each system. In contrast with our previous MSE studies of a related POSS−PU hybrid based on a diol tethered POSS16 a single T2 distribution was observed for all systems. It is thought the hard-block relaxation was too fast to be observed by the MSE measurements, and the technique therefore observes the contiguous, elastomeric phase relaxation only.

Figure 6. (a) Spin-diffusion curves plotted as the normalized intensity (I/I0) against the square root of the mixing time (tm) for all systems studied. We observe that as the % POSS inclusion in the system increases, so too does the rate of magnetization transfer from the rigid to mobile phases. Through extrapolation of the initial slope of these 1/2 are obtained which may be related the curves values of (ts,0 m) diameter of the rigid phase via eq 5. (b) Spin diffusion derived values 1/2 which are related to the effective diameter of the rigid phase of (ts,0 m) domains (drigid) by eq 5 and are tabulated in Table 2. Note that increasing levels of POSS inclusion appears to effectively decrease the size of the rigid phase domains and that there is a significant step change in this process between 4 and 6 wt % POSS, consummate with a shift in the dominant phase-domain morphology as a result of POSS inclusion in the system.

It can be observed from Figure 6a that, as the mass fraction of covalently incorporated POSS increases in the PU materials, the rates of 1H spin-diffusion from the fast-relaxing rigid domains change significantly and progressively as a function of POSS loading. If we assume an initial rate approximation,53 we can extrapolate from these data values of the rate of 1/2 magnetization exchange ((ts,0 m ) ) between the rigid and mobile phases. These magnetization exchange rates for each system are given in Figure 6b. Here, we observe that the rate of magnetization transfer from the rigid to mobile domains changes significantly as a function of POSS loading, and this is G


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Macromolecules Figure 7 shows T2 values for each derived distribution at maximum amplitude (T2Imax), plotted as a function of wt % POSS.

Figure 7. T2Imax, corresponding to the contiguous mobile phase, as a function of wt % POSS for all systems studied. Note that maximum value of the T2 distributions decrease as a function of POSS inclusion to approach a minimum at 8 wt % POSS, corresponding to a progressive decrease in proton mobility within these systems. At values above 8 wt % there is evidence of a reversal in this trend. Figure 8. Reversing (continuous) and nonreversing (dashed) signals of the PU matrix and the composites, recorded during heating.

From the T2 data shown in Figure 7 it is clear that POSS inclusion has an effect on those protons associated with the contiguous, mobile phase of the PU systems, and it becomes significantly more rigid as the levels of POSS concentrations increase toward 8 wt %. These data support the SAXS, AFM, and spin diffusion results which strongly suggest that POSS inclusion is decreasing the size of the hard-segment domains and increasing their surface area. An increased surface area, discrete rigid phase would likely lead a reduction in mobility of the surrounding elastomeric material, and these data are again also consistent with the results of our previous studies.16,54 Finally, there is also some evidence of a reversal in this stiffening trend: The rate of change in T2 decreases as it approaches a minimum at 8 wt %, and T2 actually begins to increase again at values above 8 wt % This is yet another indication of the complexity of the phase segregated domain morphology in these systems and its comparative sensitivity to the covalent inclusion of POSS into the PU matrix. Modulated Differential Scanning Calorimetry (MDSC). In agreement with the typical paradigm of polyurethanes, MDSC thermograms (Figure 8) show three distinct regions: the glass transition below 0 °C, the first hard domain relaxationoften associated with their glass transitionaround 50 °C, and, above 150 °C, a multistep microphase mixing. The glass transition step is visible in the reversing signal (−80 to 0 °C), without any indication of enthalpy relaxation in the nonreversing one. The Tg increases on the addition of POSS by up to 30 K for 10 wt % loading (Figure 9 and Table 3). The rise of Tg is significantly higher than that of the same system on inclusion of POSS as pendant groups17,18 or beads20 and is consistent with the global stiffening of the matrix as observed by MSE NMR. There are several mechanisms that may contribute to this stiffening response. The sterically large, chemical cross-links may restrict mobility either because they change the topology56 or because they immobilize chains anchored in a fashion similar to conventional nanoreinforcement. On the other hand, the change of topology also restricts microphase separationas already shown by SAXS and NMR and to be confirmed shortly by the MDSC mixing peaks. This increase in Tg may therefore

Figure 9. Several measures of glass transition temperature as recorded with different techniques against POSS content.

Table 3. Thermal Properties of Glass Transition and Microphase Mixing POSS content (wt %)

Tg (°C)

ΔcP (J/(g K))

Tmix,1 (°C)

Tmix,2 (°C)

ΔHmix,total (J/g)

0 2 4 6 8 10

−69.3 −64.8 −57.5 −38.5 −37.5 −28.6

0.30 0.32 0.33 0.40 0.41 0.49

154 162 165 168 169 172

222 212 221 223 218 216

9.7 14.5 14.3 13.7 14.0 12.7

be attributed to increased mixing of the hard segments, as described in various mixing models.57,58 We believe that this last mechanism is dominant in the studied system based on the increase of ΔcP on addition of POSS (Table 3). This trend is opposite to the one expected on immobilization due to crosslinking or anchoring59,60 but compatible with more hard segments dissolving in the soft phase and participating in the glass transition.61 In the present system, though, we may keep in mind that POSS may contribute to ΔcP, too, as their core is small as compared to the length scale of segmental dynamics, H


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Macromolecules they possess flexible organic ligands, and thus they are expected to participate in segmental mobility. Finally, we would like to emphasize that POSS addition broadens the glass transition step (Figure 8), indicating increased heterogeneity in the hybrids. Our dielectric analysis will show that this is mainly due to an inherent broadening of the α relaxation, rather than an interplay in the strengths of two different relaxations, similar to that we observed in the “beadlike” configuration.20 The first relaxation of the hard segments occurs at approximately 50 °C. It is a very weak and broad endotherm on the nonreversing signal, and thus no quantitative analysis is possible from these data. The microphase mixing occurs in two steps (MM1 and MM2, respectively) above 150 °C (Figure 8). A first peak MM1 around 160 °C on the nonreversing signal migrates to higher temperatures with increasing POSS content, indicating the presence of rather more stable structures. Its enthalpy diminishes on addition of POSS with the exception of the matrix, but please note that, in this case, a small restructuring exotherm is also involved, indicating a reduction in the quantity of the relevant structures. The opposite trends are observed for peak MM2 at ca. 220 °C. The peak temperature decreases slightly, but the enthalpy increases, eventually dominating over MM1 at 8 wt %. This trend reminds of the change of dominant peak in the complex SAXS diffractogram (Figure 3) and the emergence of spherical structures on addition of POSS, as observed by AFM. Thus, we may relate the larger structures (q ≈ 3 × 10−2 Å−1 SAXS peak) to the MDSC MM1 endotherm and the elongated hard structures in AFM. Correspondingly, we may relate the smaller structures (q ≈ 6 × 10−2 Å−1 peak in SAXS) to the MM2 endotherm and the spherical structures. The overall reduction of the microphase mixing enthalpy (Table 3) is consistent with the suppression of the SAXS diffractograms and the NMR data and confirms the decrease of microphase separation on the addition of POSS. Dynamic Mechanical Analysis (DMA). Three events occur in the DMA thermograms (Figure 10): the local β relaxation, the segmental α relaxation, and the softening of hard microdomains. We will follow them in the order of increasing temperature. A weak peak around −70 °C in the E″ thermograms for the higher loading materials reflects the β relaxation.62 This relaxation for low loading materials overlaps with the following α peak and thus is not visible. Its position does not change with loading, reflecting the localized nature of the relaxation. The α relaxation associated with the dynamic glass transition dominates the thermograms as a decreasing step in E′ and as a peak in E″ and tan δ. The peak temperature follows the trends of the calorimetric Tg (Figure 9). In agreement with the DSC steps, the DMA peaks become broader on addition of POSS reflecting the increase in heterogeneity. The Arrhenius traces (Figure 11) are slightly concave as expected for a cooperative relaxation and move toward the low frequencies (high relaxation times)−high temperature side of the graph. We will further discuss α dynamics upon presentation of data from dielectric relaxation spectroscopy. Around 50 °C, a broad peak in E″ and tan δ reflects softening in the hard domains. This peak is not discerned for high loading materials, possibly because the α peak masks it or because the effect diminishes, much like the observations by MDSC.

Figure 10. Real and imaginary parts of Young’s modulus and tan δ for all materials under investigation.

Figure 11. Arrhenius map for all materials under investigation, with data obtained by DRS and DMA. TSDC α peak temperatures are also plotted at the equivalent τ = 100 s and Tg as measured by MDSC at the frequency of modulation ( f = 1/60 Hz). Open symbols: DMA; symbols with gray interior: TSDC and MDSC. DRS symbols are filled with black and red corresponding to the α and α′ relaxations, respectively.

Despite the cross-linking, the rubbery modulus decreases monotonically, down to approximately a tenth of the value of the matrix. It appears that the hard microdomains that prevail in the low loading materials reinforce the matrix to a much greater extent than the chemical cross-links imposed by the particles. Thermally Stimulated Depolarization Currents (TSDC). TSDC thermograms reveal three distinct regions I


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reflected on the dielectric loss is severely decelerated in agreement with all previous techniques (Figure 13). Please note again that the low (≤4 wt %) and high (≥6 wt %) loading materials show significantly different spectra.

which we will follow in the order of increasing temperature (Figure 12).

Figure 12. Thermocurrents recorded with all materials under investigation. Figure 13. Dielectric loss spectra in the region of segmental relaxations.

At ca. 120 °C the β relaxation is visible for all materials. Its position does not change, in agreement with DMA results. The difference in the peak temperature is attributed to the different frequencies of the two techniques. The α peak dominates the thermograms from −80 to 0 °C. In agreement with the previous techniques, the peaks move to higher temperatures upon addition of POSS and become progressively wider. Peaks of materials with loading higher than 6 wt % are visible only as shoulders on the following, strong, MWS peak (to be described in the following). The α peak temperature is considered a good measure of the calorimetric Tg, and indeed, for the low loading materials they correspond very well (Figure 9). The Arrhenius points at 1.6 mHz also correspond well with the traces from DMA. The large peak dominating the range between −10 and 30 °C is the interfacial Maxwell−Wagner−Sillars relaxation. This corresponds to the depolarization of free charge carriers accumulated on the interfaces of areas with different conductivities. Interestingly, in this case, the peaks are structured, as most easily discerned for the high loading materials. The matrix peak is roughly in the same position as the high temperature component, while the low temperature component rises with increasing POSS content. This behavior reminds that of MDSC microphase mixing endotherms and the relevant peaks in SAXS diffractograms. The low temperature TSDC component may be related to the high temperature MDSC endotherm, the peak at higher q-values and the spherical structures observed in AFM. Dielectric Relaxation Spectroscopy (DRS). Details on the dielectric response of the PU matrix and other PU−POSS hybrids have been published in our earlier papers.17,18,63 Briefly, four molecular mobility relaxations are typically observed: (i) local γ relaxation due to crankshaft motion of methylene sequences along the macrodiol contour; (ii) local β relaxation related to carbonyl groups on the urethane bonds; (iii) the dynamic glass transition−α relaxation; (iv) a weak α′ relaxation, attributed either to restricted segmental motion of PTMEG, anchored on rigid structures,20 or local mobility of the hydrogen bonds in the hard domains.64 Upon cross-linking with octa-OH-POSS, the secondary γ and β relaxations remain unaffected, reflecting their highly localized nature, and therefore, we will not comment on them further in this article. On the other hand, the segmental mobility as

Starting from the high frequency side, a prominent peak for the matrix and the low loading materials corresponds to the α relaxation. It moves to slightly lower frequencies on addition of small amounts of POSS. The α′ for the same materials is reflected in the unreasonably low slope, or a weak shoulder on the low frequency wing of α, around 102−103 Hz. Finally, a steep slope on the low frequency side corresponds to the combined effect of dc conductivity and MWS relaxation. We now turn our attention to the high loading materials. Here, the maximum of dielectric loss occurs at much lower frequencies, in the region of 102 Hz. Judging from the Arrhenius map of β relaxation and the sequence of lower temperature spectra, a weak shoulder above 10 5 Hz corresponds to the β relaxation. However, interpretation of the spectra of high loading hybrids is not straightforward for the following reasons: (i) A standard fitting procedure with two Havriliak−Negami terms and a slope at low frequencies17,20 is inadequate and provides unreasonably high strength for the β relaxation. (ii) At higher temperatures (spectra not shown here), the segmental relaxation area, now extending 3−4 decades in frequency, does not form a peak but rather has a complex structure with a shoulder in the same region as the α′ of the matrix and the low loading materials. It is thus, obvious that there is more than one relaxation contributing to the overall dielectric response in these materials. Fitting Procedure. Following common practice, we fitted a sum of the following terms to the spectra in the region of segmental dynamics: (i) Three Cole−Cole terms ⎛ ⎞ Δε ⎟ ε″(f ) = Im⎜⎜ a⎟ ⎝ 1 + (if /fmax ) ⎠

(6)

to account for the three relaxations (α, α′, β). In this model f is the field frequency, f max the characteristic frequency of each relaxation, Δε the contribution of each relaxation to the total dielectric constant, i.e. the strength, and a an exponent describing the width of the relaxation, with a = 1 corresponding to the Debye single relaxation time peak. Although segmental relaxations are known to follow rather the asymmetric J


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over the enthalpic one in Gibbs free energy, and the behavior is typical in glass forming liquids and linear polymers.33,65,67 An increase is interpreted as a gradual release of constraints imposed by physical cross-links, like crystallites68,69 or strongly interacting particles.70−72 Here, the stability for low loadings may be interpreted as a compromise between the two phenomena, i.e., release of constraints imposed by the hard nanodomains and interplay of the two terms in the Gibbs free energy. Upon addition of POSS, despite the chemical cross-linking, the decreasing trend in Δεα with increasing temperature intensifies. This may be interpreted as a decrease in the effective cross-linking density, in agreement with our previous comment that hard nanodomains are more effective cross-links than the octafunctional POSS. We note also that, at relatively low temperatures (ca. 0 °C), Δεα increases systematically with POSS content, up to a factor of 10. Apparently, this happens due to the inhibition of microphase separation and the introduction of POSS into the soft phase, where more material now contributes to the glass transition. This observation and reasoning are in agreement with the increase of ΔcP measured by MDSC. Shape. On incorporation of POSS cross-links, the relaxation widths increase, as manifested by the lower aα values (Figure 15). This is in agreement with earlier observations on cross-

Havriliak−Negami model, the overlap of the peaks does not allow for the calculation of asymmetries. (ii) A log−log slope ε″ = Af s to account for the charge mobility contribution. In order to further reduce the fitted parameters, we also fixed some of them that did not exhibit significant variations in temperature regions where they were reliably calculated. Namely, (i) the strength Δε and width a of the β relaxation to the values calculated at −45 °C (the highest temperature where it is clearly visible for all samples), (ii) the slope exponent, s, to the value calculated at 55 °C, and (iii) the width α of the α′ relaxation at the value calculated in the middle of the temperature region where the most unambiguous fit was achieved. Thus, we quantified, for the α and α′, the time scale in terms of f max, the strength in terms of Δε, and width in terms of a, and in this order we will present them in the following. Time Scale. Interestingly, the time scale of α′ is the same for both the low- and high-loading materials (Figure 11), indicating that they are of the same nature, and justifying assignment of the same symbol. It is not clear whether the temperature dependence is linear (Arrhenius) or slightly concave (very strong VFT, where strong used here is the opposite of f ragile65). The cross-linking by POSS significantly decelerates the α relaxation in agreement with DMA, TSDC, NMR, and DSC. Please note here that the DMA trace is shifted to higher frequencies−lower temperatures. This has been also observed by Santangelo and Roland in polyisoprene and happens because retardation times are always longer than the corresponding relaxation times.65,66 For high loading materials, the two techniques differ also in the concavity, presumably because α′ has a contribution to the DMA peaks. We would like to remind here that the DMA peaks are actually a lot broader for these samples, with indications of a double structure. Relaxation Strength. The temperature dependence of the strength parameters confirms the different structures of the high and low loading materials (Figure 14). The matrix and the

Figure 15. Shape exponent aα of α relaxation as a function of temperature. Low values correspond to broader peaks.

linked systems73 and interpreted as a broader distribution of relaxation times. This is typically attributed to a higher spatial inhomogeneity.72 With increasing temperature, the relaxation peaks become progressively narrower, which is a typical behavior of α relaxation. The effect is more prominent with increasing POSS content. Again, this trend may reflect homogenization of the system.

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CONCLUSIONS A model polyurethane system has been cross-linked with octafunctional POSS moieties, and the resulting complex organic−inorganic hybrid networks have been investigated with respect to their morphology and molecular mobility. Despite their crystalline nature, the POSS units disperse in the nanoscale, indicating successful chemical binding on the polymer chains. With increasing loading, more diisocyanate moieties are attached to the reactive ligands of the particles, become unavailable for the formation of hard structures, and thus, the degree of microphase separation is highly suppressed.

Figure 14. Strength parameters vs temperature for both relaxations.

2% hybrid show rather stable Δεα(T) and Δεα′(T) . On the other hand, starting from 4%, Δεα(T) shows a decreasing trend, partly compensated by an increase in Δεα′. This dependence becomes stronger upon incorporation of POSS. Please note, though, that the decrease of Δεα(T) is higher than the increase of Δεα′ by a factor of 10. A decrease in Δεα with increasing temperature is usually interpreted in terms of gradual dominance of the entropic term K


Macromolecules

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The hard domains tend to shrink in size and total volume. Two kinds of structures seem to occuran “elongated” and a “spherical”with one gradually prevailing over the other with increasing loading. POSS does not appear to affect intersegmental distances in the hard domains, and thus we believe that they reside in the soft phase. NMR, MDSC, DMA, TSDC, and DRS, in agreement with each other, confirm that the segmental dynamics of the soft phase slows down as a result of crosslinking and phase mixing, but interestingly, this happens in a steplike manner between 4 and 6 wt % loading. This region corresponds also to steplike changes of the size of hard domains and the distances between hard segments dissolved in the soft phase. It also corresponds roughly to the prevailing of “spherical” type of hard domains over elongated hard structures. An α′ relaxation, slower than the main dynamic glass transition (α relaxation), also contributes to segmental dynamics, as studied by thermal, thermomechanical, and dielectric methods. Its dynamics is not altered by addition of POSS. Strikingly, despite extensive chemical cross-linking, the rubbery mechanical modulus decreases on addition of POSS by an order of magnitude, confirming that hard microdomains reinforce the matrix to a greater extent than chemical crosslinks. With this article, we conclude a series of investigations on PU−POSS elastomers with particles embedded in the soft matrix with different binding modes. What this study, alongside with our previous work, has shown is that the effects and influence of POSS on polyurethanes are complex, multiscaled, and often counterintuitive. For at least segmented polyurethanes (and contrary to many expectations), POSS is disruptive rather than reinforcing, and while it may improve some aspects of thermal stability, it degrades the microphase segregation of polyurethane elastomers which is the source of many of their desirable physical properties. In future work, it will be informative to follow effects of POSS which are not chemically bound on the PU matrix but rather mixed with it in a conventional nanocomposite configuration. Investigations of the effects of POSS inclusion on the morphology and dynamics of systems that are less prone to microphase separation (i.e., systems with aliphatic instead of aromatic diisocyanates) may provide additional insight into the behavior of POSS−polymer hybrid organic−inorganic networks.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (K.N.R.). Present Address

K.N.R.: Technische Universität München, Physik-Department, Fachgebiet Physik weicher Materie, James-Franck-Str. 1, 85748 Garching, Germany. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge technical assistance with DMA measurements by Dr. Joanna Pagacz (CUT). The analysis of dielectric and X-ray data was performed with the software graf ity, created and maintained by Dr. Daniel Fragiadakis. This work has been cofunded by the National Science Center in Poland under Contract No. DEC-2011/02/ A/ST8/00409 (K.N.R., E.H., M.J., and K.P.). This research has been cofinanced by the European Union (European Social Fund ESF) and Greek national funds through the Operational Program Education and Lifelong Learning, Research Funding Programs Aristeia and Thales (S.K. and P.P.). Portions of this work were performed under the auspices of the U.S. Department of Energy by Lawrence Livermore National Laboratory under Contract DE-AC52-07NA27344 (J.P.L., H.E.M., and S.J.H.). The final version of this manuscript was formed during S.K.’s COST MP1105-17395 short term scientific mission (STSM) to CUT.

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S Supporting Information *

Figure S1: AFM phase images illustrating certain features of the materials (a, higher magnification image on the border of a globule in the matrix, illustrating the difference in microphase separation; b, irregular particles dispersed in the continuous phase of the 4 wt % hybrid; c, closeup on a set ofpresumably POSS containing“sacs” in the 10 wt %; d, the same sacs at higher magnification); Figure S2: deconvolution of the WAXS diffractogram of the PU + 4 wt % POSS hybrid to its constituent curves; Figure S3: full width at half-maximum of the X-ray peaks in the high q region. This material is available free of charge via the Internet at http://pubs.acs.org. L


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