POSS along the Hard Segments of Polyurethane. Phase Separation

Article pubs.acs.org/Macromolecules

POSS along the Hard Segments of Polyurethane. Phase Separation and Molecular Dynamics Konstantinos N. Raftopoulos,*,†,‡ Małgorzata Jancia,† Dionysia Aravopoulou,‡ Edyta Hebda,† 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 S Supporting Information *

ABSTRACT: We report for the first time on a polyurethane system with polyhedral oligomeric silsesquioxanne (POSS) particles chemically bonded along its main chain as nanobuilding blocks. Morphology and molecular dynamics of this novel system are studied with imaging (SEM), thermal (DSC, MDSC) dielectric (TSDC, DRS), and mechanical (DMA) techniques. Two distinct phases occur: one chain extended with POSS and one with chemistry and morphology similar to the matrix. Although Tg remains practically unaffected upon incorporation of particles, a relaxation α′ slower than the main dynamic glass transition α and present also in the matrix is enhanced in the presence of POSS. We attribute it to decelerated dynamics in the presence of heavy moieties like particles and rigid hard microdomains. No effect on the time scale of these relaxations is imposed by the presence of POSS, but rubbery modulus is significantly reduced due to the soft POSS-extended phase.

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INTRODUCTION The dispersion of reinforcing nanoparticles in polymer matrices and their targeted spatial placing are key concerns of the design and synthesis of novel polymeric nanomaterials. Both are facilitated by the “nanobuilding block” approach, by which the particle is properly modified in order to be covalently incorporated in the macromolecular structure itself.1 Polyhedral oligomeric silsesquioxannes (POSS) have been proven to be ideal candidates for the implementation of this approach.2−4 They consist of polyhedral cages with Si atoms on the vertices, interconnected with Si−O−Si linkages forming the edges. One of the four Si substituents, known as the “vertex group”, does not participate in the cage formation and may be chosen practically between all known organic groups, reactive or not, in order to facilitate solubility in the matrix or chemical bonding in desired locations on the macromolecular chain. Depending on the number of vertex groups that possess a reactive moiety, several chain topologies are possible. Particles without reactive vertex groups will be blended in the polymer and form conventional nanocomposites. Particles with one reactive group will be tethered as side chains or as end-caps, two will result in a “bead structure” with POSS lying along the main chain, and particles with more than two reactive vertex groups will typically act as “heavy” chemical cross-links. Nanoparticles are known to affect segmental dynamics by two direct mechanisms. Direct POSS−polymer interactions immobilize part of the polymer and increase Tg as shown for PDMS/silica5 and PMMA/silica6 nanocomposites. Increase of free volume facilitates segmental mobility and lowers Tg, due to © XXXX American Chemical Society

loosened molecular packing of the chains, as shown by dielectric methods in polystyrene/phenethyl-POSS nanocomposites in a work by Hao et al.7 On the other hand, indirect effects may occur especially if chemical reactions proceed in the presence of particles and even more if the particles are involved in the chemical process. To mention a few, particles may impose changes in molecular weight,8 crystallinity,9 or microphase separation.10 In the case of POSS, the final effect on segmental dynamics depends on the polymer matrix, the nature of their vertex groups which determine the particle-chain interactions and the volume of the particle, and finally the network topology (ref 4 and references therein). The final change of Tg at a typical content of 10 wt % is known to vary from approximately −25 °C for isobutyl-substituted POSS tethered on isobornyl methacrylate11 to +80 °C for POSScross-linked polyimide.12 Polyurethanes (PU) are essentially copolymers of flexible macrodiols (soft segments) and sequences of diisocyanates and short diols (hard segments). Microphase separation is the characteristic which determines most of their physical properties.13 It is driven mainly by the chemical incompatibility between the hard and soft segments of the polymer and the formation of hydrogen linkages between the urethane bonds on the former.14 In PU, prediction of the effect of modification is more complicated as the particles are expected to alter the Received: July 6, 2013 Revised: August 24, 2013

A

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Scheme 1. Synthesis Route for the PU−POSS Hybrids

degree of microphase separation (DMS), which is known to have a major impact on thermomechanical properties.13 A decrease of DMS is expected to increase Tg as more, rigid, hard segments are diluted in the soft phase.15 In PU systems with POSS tethered on the hard domains as pendent groups the general trend is a moderate increase of Tg.16−23 The magnitude of the effect depends on the components of the PU and the bonds employed for the tethering. Tethered POSS also tend to inhibit microphase separation.23−25 In a previous article on PU systems with POSS tethered as side groups to the hard domains, we have reported that direct particle−chain interaction is the dominant mechanism increasing the Tg when the length of the PU segments is large, while decrease of microphase separation becomes more important when the segments are shorter.23 In this article, we continue our study of PU−POSS materials with a system where POSS is incorporated along the chain contour of a common PU and more specifically on its hard segments, acting as chain extender. Organic−inorganic hybrid materials with POSS incorporated along the chain backbone do not appear often in the literature, and to the best of our knowledge, no such polyurethane system has been reported up to now. In this article, after having prepared such a PU−POSS system in a two-step synthesis

procedure, we investigate molecular dynamics and to some extent phase separation. The research is conducted by means of imaging (scanning electron microscopy, SEM), thermal (conventional and modulated differential scanning calorimetry, DSC, MDSC), dielectric (thermally stimulated depolarization currents, TSDC, and dielectric relaxation spectroscopy, DRS), and thermomechanical (dynamic mechanical analysis, DMA) methods. We will show that while Tg remains practically unaffected, a molecular mobility mechanism a little slower than the dynamic glass transition is significantly enhanced by the addition of POSS. This mechanism is probably related to dynamics of macrodiol chains anchored to rigid structures.

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EXPERIMENTAL SECTION

Materials. The polyurethane matrix was prepared by reaction of 4,4′-diphenylmethane diisocyanate (MDI, Aldrich) as isocyanate component and poly(tetramethylene glycol) (Terathane 1400, Invista) with molecular weight of ∼1400 as the elastic component. 1,4Butanediol (BD, Aldrich) was used as chain extender. Nanocomposites were prepared by appropriate substitution of chain extender by disilanol isobutyl POSS (Hybrid Plastics, DSIPOSS) in order to provide samples with POSS mass fractions of 2−10%. Structure of the particles is shown in Scheme 1. The mass fraction of elastic component in the polyurethane is always 50%. B

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Figure 1. Top: backscattered electron images of the matrix and the hybrids with 4 and 10 wt % loadings. Scale bars are 50 μm. Bottom: qualitative EDS spectra recorded at points noted on the micrographs. where A is the area of the plates, d the sample thickness, and V the polarizing voltage. The specimens in this investigation are films of thickness ca. 1 mm 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. Dielectric Relaxation Spectroscopy. Standard dielectric relaxation spectroscopy (DRS) measurements29,30 were performed on the same specimens as with TSDC in the region 10−1−106 Hz with a Novocontrol Alpha analyzer and in the temperature region −130 to 110 °C. For a few selected samples, the temperature range was extended to 230 °C in order to follow charge mobility effects. Dynamic Mechanical Analysis. DMA thermograms were recorded with a Netzsch 242C dynamic mechanical analyzer in the temperature range −110 to 120 °C at 2 K/min. Frequencies 1−50 Hz were used. Samples were beams of approximate size 22 × 7 × 2 × nm, measured in a three-point bending configuration on a 10 mm sample holder.

Before use, PTMEG, BD, and DSIPOSS were dried overnight under vacuum; MDI was used as received. MDI was charged into a 100 mL three-necked round bottomed reactor, equipped with a mechanical stirrer, a thermometer with a temperature controller, and an argon inlet. The MDI was heated to 70 °C, and a solution of DSIPOSS in a suitable amount of terathane polyol was then added in one portion. The polymerization reaction was performed under an argon atmosphere at 80 °C for 2 h to form a polyurethane prepolymer. The NCO group content was then determined, and the prepolymer was mixed with a suitable amount of 1,4-butanediol. The resulting mixture was poured out on a Petri dish, cured at 110 °C for 2 h, and finally postcured at 80 °C for 16 h to form a solid elastomer. The synthesis procedure is illustrated in Scheme 1. Scanning Electron Microscopy. Backscattered electron images of the matrix and selected hybrids were recorded by a JEOL JSM-6010LA scanning electron microscope with energy dispersive X-ray spectroscopy (EDS) function. Uncoated cryofractured surfaces were imaged at low vacuum mode (60 Pa). Differential Scanning Calorimetry. Conventional differential scanning calorimetry (DSC) experiments were performed in the region −120 to 20 °C at rate 10 K/min, following a cooling at the same rate from ambient temperature. Modulated differential scanning calorimetry (MDSC) experiments were performed from ambient temperature up to 350 °C at rate 5 K/min and modulation 2 K over 60 s. For both DSC and MDSC a TA Q200 calorimeter with Tzero technology26 cooled with liquid nitrogen and purged with nitrogen was used. Specimens had mass 7.0−8.5 mg and were placed in standard aluminum pans. Thermally Stimulated Depolarization Currents (TSDC). Thermally stimulated depolarization currents (TSDC) is a special dielectric technique in the temperature domain. It roughly corresponds to measuring dielectric loss in the region of mHz,27,28 thus extending frequency range to low frequencies, typically not accessible by conventional dielectric techniques. By 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 shortcircuited 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 as28

Inorm =

I Id = EA VA

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RESULTS AND DISCUSSION Scanning Electron Microscopy. All materials have a network of interconnected pores, which with increasing loading becomes denser and with more irregular geometry as evidenced at low magnification (not shown here). Images inside those pores are presented in Figure 1. The matrix is composed by spherical structures of a few micrometers. These structures have been known to consist of a dendritic backbone of tangentially arranged, crystallized hard segments, with soft segments filling the area between the branches.31,32 Addition of 4 wt % POSS gives rise to one more phase of micrometer-sized lamellae, radially arranged in loose spherulites. This phase lies on both the pores and the fractured surface, but more regular arrangement is possible in the former. As proven by the intense Si peak in the respective EDS spectra (Figure 1, bottom row), this phase is richer in POSS particles than the “regular” grainy area. Nevertheless, it is not a purePOSS phase because it (i) is very unstable and easily molten by the 10 keV beam, after a few scans, and (ii) differs in morphology and EDS spectrum from a pure POSS crystal which may be found in the Supporting Information (Figure S1). Since chain extension by BD was held at a different time than the extension by the particles, we expect that two polyurethane phases are formedone POSS-extended and one BD-extendedwhich are partially immiscible to each other.

(1) C

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Moreover, it is expected that the urethane groups of the first phase are sterically hindered by the adjacent bulky particles and therefore unable to form hydrogen bonds and eventually hard microdomains. A similar phase separation has been reported in blends of poly(ethylene oxide) (PEO) grafted on POSS particles and unmodified PEO.33 At 10 wt % loading, the “grainy” PU structure is completely disrupted, but two phases are still visible as dark areas and bright lamellae, both inside the pores and at the bulk (fractured surface). At this high loading, some structures resembling purePOSS crystals are also detected especially in the rim of pores. Thermal Analysis. Soft Phase Glass Transition. In the subambient temperature region, the glass transition step of all materials is clearly observed in the DSC thermograms (not shown here). Tg calculated as fictive temperature34 is plotted against POSS content in Figure 2 together with onset and end

Figure 3. Reversing and nonreversing heat flow thermograms recorded with all materials under investigation during heating. Lines have been translated vertically for clarity.

the typical polyurethane paradigm. Following the nonreversing signals, the first weak and broad endotherm below 100 °C has been attributed to to disruption of ordering of short MDI-BD sequences35,36 or glass transition of hard microdomains.37,38 The next endotherm around 170 °C reflects the main phase separation.15 The last peak, over 250 °C, is attributed to the eventual decomposition of the materials, as it is accompanied by a sharp drop in the reversing signal and occurs at temperatures close to those observed for similar systems.39,23 Interestingly enough, no melting peak of the POSS extended phase is observed albeit SEM shows well-defined structures. This endotherm should arise above room temperature, and below the melting peak of POSS, which according to observations in polarizing microscope is around 85 °C. However, we would like to note that SEM specimens, before measurement, have been exposed to low temperatures during cryofracturing while DSC ones did not and therefore supercooling might not be enough to allow for nucleation. The second endotherm is accompanied by a weak step in the reversing signal, reflecting that more material now contributes to the total specific heat. Its peak temperature Tmix, along with the enthalpy ΔHmix are tabulated in Table 1. Tmix does not show any monotonous trend with loading. ΔHmix on the other hand is decreasing with addition of POSS, albeit not monotonously. This overall trend might reflect the gradual dominance of the POSS-extended phase over the BD-extended. We now turn our attention to the decomposition peak above 250 °C. Its onset is accompanied by a sudden drop in the reversing signal reflecting the decrease of heat capacity as a result of the change in the chemical nature of the material and mass loss due to gaseous products of the decomposition. While the neat PU decomposes at a single step, the thermograms of the hybrids show a multistep process, reflecting the plurality of phases and chemical bonds present in them. The chemical procedure associated with each step merits a detailed investigation by analytical methods in future work. Dielectric Analysis. Thermally Stimulated Depolarization Currents (TSDC). In Figure 4 we show the normalized TSDC thermograms with all materials under investigation. In the studied region, four relaxations are observed. Starting from low

Figure 2. Calorimetric glass transition temperature Tg by DSC, α relaxation peak temperatures Ta by TSDC, and TE by DMA at 5 Hz, plotted against POSS content. Dashed lines are onset and end temperatures of the step in DSC thermograms and the 90% limits of the relaxation peaks in TSDC experiments.

Table 1. Key Thermal Properties As Measured by Thermal Techniques POSS (wt %)

Tg (°C)

Δcp (J/(g K))

Tmix (°C)

ΔHmix (J/g)

0 2 4 6 8 10

−53.1 −54.7 −54.3 −52.3 −52.9 −52.4

0.48 0.44 0.45 0.47 0.48 0.47

167.2 166.0 162.5 169.7 174.8 161.5

20.0 25.0 15.8 19.8 16.7 9.3

temperatures (black dashed lines) and tabulated in Table 1. No significant change of Tg with POSS content is observed, indicating that incorporation of POSS does not alter significantly the mobility of the soft phase. However, the end temperature of the glass transition step is increased for loadings higher than 4 wt %. We will come back to this point later and comment on it in light of the results of dielectric analysis. Interestingly, the heat capacity step Δcp (Table 1) shows no monotonous dependence on POSS content and lies in the region 0.44−0.48 J/(g K). Microphase Mixing and Decomposition. Microphase mixing was studied by means of modulated differential scanning calorimetry. Reversing and nonreversing heat flow thermograms in the region 50−340 °C are shown in Figure 3. The thermograms exhibit three distinct regions, in agreement with D

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Figure 4. TSDC thermograms recorded with all materials under investigation.

Figure 5. Comparative dielectric loss spectra at −15 °C. Representative fitting curves for the 6 wt % hybrid are also shown.

strongly local character, in consistency with observations in ref 20. Therefore, we will not comment further in this article. Segmental Dynamics. In Figure 5, we show the dielectric loss spectra of the PU and the hybrids at −15 °C. It is clear that the relevant dominance of α′ is enhanced with incorporation of POSS, but the strong overlap between the two relaxations does not allow for safe conclusions regarding variations in time scale. In order to clarify this and to quantify the relevant dominance of the two relaxations, we performed a typical fitting procedure of model functions to the experimental data. An example of the analysis is included in Figure 5. Because of the strong overlap, several assumptions were made in order to reduce the number of free parameters and eventually obtain meaningful results:42,43 Although peaks related to segmental dynamics are known to be asymmetric in shape,29 symmetric Cole−Cole terms were used to model each of the peaks.43 A term of the form ε″(f) = Af−n was also added to account for dc conductivity and other possible relaxations, resulting in a strong negative slope at low frequencies for high temperatures. The exponent n was determined at a high temperature, typically +5 °C, where the slope was clearly visible and then was fixed. Eventually, fitting was performed on the dielectric loss (ε″) spectra according to the equation

temperatures, at −117 °C the β relaxation peak reflects motion of carbonyl groups of the urethane bond, probed by attached water molecules.20,40,41 In agreement with results in our earlier work,20 β is not affected by addition of particles as a result of its very localized nature. The peak around −60 °C corresponds to the dynamic glass transition (α relaxation). Its peak temperature, Tα, is a good measure of the calorimetric Tg.27 In Figure 2, we show Tα dependence on the POSS content as gray circles. Strikingly, Tα is not independent of POSS content like Tg but increases by several degrees after addition of more than 4 wt % POSS. A closer look in Figure 4 shows that the α peaks of the hybrids with POSS content ≥6 wt % are significantly broader than their lower content counterparts, but only at the high temperature side. We have quantified this effect by measuring the temperatures where the normalized current has 90% of its maximum value for each one of the peaks and depicting them as gray lines in Figure 2. The observation is in agreement with the broadening of the glass transition step in DSC thermograms. It is not clear at this point whether the increase of Tα and the broadening are due to effects on the α relaxation or due to the rise of a second, high temperature component, overlapping with the α peak. The answer will be given, upon evaluation of the data from dielectric relaxation spectroscopy, in the next section. At temperatures above −20 °C, Maxwell−Wagner−Sillars (MWS) relaxation becomes dominant. This interfacial relaxation reflects the mobility of charge carriers trapped at the interfaces of regions with different conductivity and is characteristic of inhomogeneous materials. We will not further comment on those peaks at this point as most of them occur in the region above the polarization temperature. Dielectric Relaxation Spectroscopy (DRS). Dielectric loss spectra in a wide temperature range are qualitatively similar to those we have published elsewhere20 as well as included in the Supporting Information (Figure S2). Briefly, in the −150 to 100 °C region, four peaks are visible. In addition to β and α peaks already seen by TSDC, a local γ relaxation faster than β is present at low temperatures, while a shoulder on the low frequency side of the α peak (Figure 5) is designated as α′. At higher temperatures, the low frequencies side of the ε″ spectra are dominated by a steep slope due to conductivity related phenomena, i.e., the dc conductivity and the interfacial MWS relaxation. Comparative plots in the sub-Tg temperature region (not shown here) show that time scale of secondary relaxations is not affected by the presence of nanoparticles as a result of their

⎛ ⎞ ⎛ ⎞ ⎜ ⎟ ⎜ ⎟ Δε Δε′ ⎜ ⎟ + Af −n ⎜ ⎟ ε″(f ) = Im aCC + Im CC ⎟ a ′ ⎜ ⎛ ⎞ f ⎜⎜ ⎟ 1 + ⎜i f ⎟ ⎟ ⎜1 + i f ⎟ ⎝ 0⎠ ⎠ ⎝ f ′0 ⎝ ⎠

( )

(2)

The first two terms are the imaginary part of the Cole−Cole model for α and α′ relaxations, respectively. In this model Δε is the intensity of the relaxation, i.e., its contribution to the dielectric constant of the material. f 0 = 2π/τ0 is the field frequency corresponding to the characteristic relaxation time τ0. The exponent αCC is a measure of the width of the peak and subsequently the distribution times, aCC = 1 corresponds to the single relaxation time relaxation (Debye relaxation), and aCC decreases toward 0 as the distribution becomes broader. Primed symbols are used to denote the same parameters for the α′ relaxation in the second term. The last term corresponds to the negative slope at low frequencies as described earlier. Since this slope is a combination of phenomena, its parameters A and n have no unambiguous physical interpretation. The results on time scale are presented in the Arrhenius map of Figure 6. α relaxation traces are concave, as expected by the Vogel−Fulcher−Tammann−Hesse (VFTH) model, valid for cooperative processes:44 E

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Table 2. Relative Dominance of α′ and Shape Characteristics of α and α′ Relaxations at −15 °C Δε′/Δεtot

aCC

a′CC

0 2 4 6 8 10

0.20 0.33 0.32 0.43 0.56 0.54

0.27 0.26 0.27 0.25 0.27 0.25

0.35 0.25 0.25 0.26 0.24 0.24

environment of α′-dipoles is evident upon incorporation of particles. Charge Mobility. At high temperatures, charge mobility dominates the spectra. Macro- and mesoscopic movement of carriers through the bulk of the material is influenced by its microstructure. Therefore, charge mobility may provide information on the morphology. It is often convenient to study free carrier phenomena in terms of complex conductivity σ* or complex electric modulus M*.46,47 M* is the inverse of dielectric function:

Figure 6. Arrhenius maps of all materials under investigation in the area of segmental relaxations. DRS points are results of fitting. DMA points are peak temperatures of loss modulus E″. TSDC peak temperatures have been plotted at the equivalent frequency of 1.6 mHz (τ = 100 s). Dashed lines are fits of the VFTH equation to the data from all materials.

fmax = f0 e−B /(T − T0)

POSS (wt %)

(3)

in which f 0, B, and T0 (Vogel temperature) are parameters of the material. We cannot make a conclusive argument on the curvature of α′ traces which would allow for a definite answer whether it is a cooperative relaxation (if concave) or not (if linear). For both relaxations, due to scattering of the points, no effect of POSS on the time scale of either relaxation is justified. Therefore, the VTFH model was fitted for each relaxation to data from all samples collectively, in order to improve statistical significance.45 The fitted lines are included in Figure 6. The obtained parameters are for α: log( f 0/Hz) = 9.7, B = 874 K, T0 = 185 K and for α′: log( f 0/Hz) = 12, B = 4033 K, T0 = 110 K. On the same plot, inverse Tα is plotted at the equivalent frequency of τ = 100 s.27 Interestingly enough, the VFTH projection of the α relaxation traces is in good agreement with the TSDC point of the matrix for low content hybrids where α is dominant. On the other hand, α′ traces are projected on the TSDC line at higher temperatures than those of α and closer to the Tα of the high loading materials. These observations, together with the gradual enhancement of α′, explain the broadening of the TSDC α peak for hybrids with more than 6 wt % POSS. In TSDC α′ is not discerned as a stand-alone peak because at this low equivalent frequency, it approaches and overlaps with α. Moreover, now it is confirmed that increase of peak temperature Tα in TSDC is a result of the relative strengthening of α′ rather than deceleration of α. We now turn our attention to the strength of the relaxations. Porosity may influence Δε and Δε′ and following the variation of those values with POSS content may not reflect well the effects of POSS in the nanoscale. However, the ratio Δε′/Δεtot = Δε′/(Δε + Δε′) reflects well the relative contribution of α′ to the total strength and is included in Table 2. We observe that, starting from 20% for the neat PU, at high contents α′ accounts for more than half of the segmental dynamics contribution to the static dielectric constant. The shape exponents αCC and α′CC of α and α′, respectively, are also presented in Table 2. A lower value of these exponent corresponds to broader peaks and, indirectly, to more heterogeneous environments for the dipoles, in the nanoscale. We may see that the environment of α-dipoles is relatively unaffected while a significant increase of heterogeneity in the

M* =

ε′ ε″ 1 = 2 +i 2 = M′ + iM″ 2 ε* ε′ + ε″ ε′ + ε″2

(4)

The typical paradigm of imaginary modulus (M″( f)) spectra of polyurethanes at high temperatures consists of two peaks.48 One of them is the Maxwell−Wagner−Sillars (MWS) interfacial relaxation due to accumulation of charges between hard and soft phases. The other is the so-called conductivity relaxation (CR) and reflects the transition of conductivity from frequency independent (dc) at low frequencies to frequency dependent (ac) at high frequencies.48,49 Below its peak frequency f CR the macroscopic transport of carriers is dominant, whereas at higher frequencies only localized motion is possible. In Figure 7, we have plotted σ′ and M″ spectra at an elevated temperature, yet lower than the onset of microphase mixing

Figure 7. Frequency dependence of conductivity (top) and imaginary electric modulus (bottom) at 130 °C for the matrix and selected hybrids.

endotherms. Here, the dominant structure in M″(f) is a rather broad peak in the 10 Hz region, designated as C1. For the pure matrix an additional peak, C2, is visible as a shoulder on the high frequency side of C1. In the σ′( f) spectra, a dc conductivity σdc plateau is followed by a step-like structure around 1 Hz corresponding to the C1 peak. A more pronounced step, around 102 Hz, only for the matrix, corresponds to the C2 peak. F

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Table 3. Activation Energies and Shape Exponents for Charge Mobility Phenomena σdc matrix 4 wt % 10 wt %

C1

C2

Eact [eV]

Eact [eV]

aCC (120 °C)

aCC (200 °C)

Eact [eV]

aCC (120 °C)

aCC (200 °C)

0.69 0.72 0.65

0.72 0.80 0.68

0.80 0.89 0.94

0.99 0.96 0.99

0.68

0.85

0.89

Upon microphase mixing, at high temperatures, f CR and σdc traces are concave, indicating that now charge mobility is connected to the molecular dynamics. In the mixed state, 4 wt % of particles does not alter a lot the charge mobility, but 10 wt % significantly suppresses it as indicated by the reduction of σdc and f CR by an order of magnitude. In a further analysis step, activation energies Eact for all traces were calculated by appropriate fitting of the Arrhenius equation

In order to clarify the origin of C1 and C2, we followed their shape and time scale with increasing temperature, with an analysis process similar to that used for the ε″(f) spectra. A Cole−Cole-like term was applied for each of the peaks. The resulting shape exponents at two different temperatures are tabulated in Table 3, and the time scale is plotted in the Arrhenius map of Figure 8. In the same figure, log σdc vs inverse temperature is also included for comparison.

fmax = f0 e−Eact / kT

or

σdc = σ0e−Eact / kT

(5)

to the low temperature data, i.e., before Tmix, and are presented in Table 3. As expected, Eact values for σdc and C1 are similar and observe the same trends. The 4 wt % hybrid has higher values than the matrix and the 10 wt % hybrid. This may indicate that POSS and POSS-extended phase at high loadings aggregate and charge carriers move through a percolated “pure PU-like” environment, but at low loadings dissolved particles inhibit the charge mobility. We would like to close this section with a few comments on the C2 peak. Its Arrhenius trace follows the same trends with C1 and σdc and has similar Eact. Therefore, we think that it is a conductivity-related phenomenon. It should be attributed to space charge polarization, at interfaces that vanish or are significantly altered upon addition of POSS. Those interfaces should be connected also to a weak endotherm observed in the MDSC thermograms of the matrix around 230 °C (Figure 3). Another footprint of C2 is hardly discerned as a weak shoulder on the low temperature side of the TSDC MWS peak of the matrix (∼−20 °C, Figure 4). Dynamic Mechanical Analysis (DMA). Storage E′ and loss E″ moduli thermograms are shown in Figure 9. Starting

Figure 8. Arrhenius map of conductivity (top) and the peaks of electric modulus (bottom) for the matrix and the hybrids with 4 and 10 wt % loading. Onset temperatures of phase mixing as measured by MDSC are shown by arrows.

Both C1 and C2 are visible even after Tmix. Additionally, only one peak is visible for the hybrids, although both MWS and CR are expected below Tmix. However, we may observe that C1 is significantly narrowed down upon mixing, as demonstrated by the difference in αCC at 120 and 200 °C, eventually becoming almost Debye with αCC,200 °C > 0.95 (Table 3). For the detailed αCC(T) dependence, please refer to Figure S3 in the Supporting Information. We think that before mixing, C1 is a composite peak with contributions from MWS and CR. Upon mixing, there are no interfaces and thus MWS vanishes. The CR, which is known to be narrow,47 remains then the only component of C1. This morphology and evolution of M″(f) peaks are very similar to those observed during melting of PTMEG.50 In support of our explanation, there is a correlation between ΔHmix (MDSC) and αCC of C1 below Tmix (Tables 1 and 3). When POSS is added, microphase separation is reduced (lower ΔHmix) and MWS contribution is weaker as reflected on the higher αCC values (narrower relaxations). f CR is known to be proportional to σdc.49 This is nicely demonstrated here in the Arrhenius maps of Figure 8. Conductivity values and f C1 follow the same trends with remarkable accuracy. For all materials, both σdc and f C1 exhibit two distinct regions separated by the microphase mixing onset temperatures, annotated with arrows in Figure 8. For the phase-separated materials, at low temperatures, traces are linear, indicating decoupling of the charge mobility from the cooperative segmental dynamics.

Figure 9. DMA thermograms at frequency 5 Hz, recorded with all materials under investigation.

from low temperatures, glass transition is visible in the area of −50 °C as a gentle step in E′ and a broad peak in E″. A little over 50 °C a shoulder in E″ accompanied by a small step in E′ corresponds to the first softening of hard microdomains observed as a weak broad peak in MDSC. G

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when it is crystallized,20,50 so it cannot be a local relaxation because then it would be visible also in the amorphous state. Therefore, our results support that α′ is related to decelerated segmental dynamics probed by terminal groups at interfaces. This is similar to the dynamics of the so-called rigid amorphous fraction (RAF) which is clearly visible in DRS spectra5 as a stand-alone peak, but in thermal experiments it is only indirectly observed as a deficit in Δcp.6 Interestingly enough, in the case of aliphatic macrodiols the slow dynamics are very similar, regardless of the structure that slowed them down, i.e., crystallites, hard microdomains, or nanoparticles.

The rubbery modulus at high temperatures is clearly reduced upon addition of particles by a factor of 3. The glassy modulus varies with loading in a nonmonotonous manner in the range 900−1300 MPa, i.e., less prominently than the rubbery; however, hybrids seem to exhibit higher values than in the matrix, similarly to the case of pendent particles.18 In the glassy state, response to the external stimulus is dependent on the dynamics which is reinforced by the particles, but in the rubbery state, toughness is controlled by the density of the physical cross-links which are absent in the POSS-extended phase as we have hypothesized earlier. Tg, as quantified by the peak temperature of E″ at 5 Hz, does not exhibit any variation in agreement with DSC observation (Figure 2). Moreover, indication of a slower relaxation is present only as a very weak shoulder in the high temperature side of the α peak for the 10 wt % hybrid. Thermal, dielectric, and mechanical methods probe different aspects of molecular mobility, and we may assume that α′ has very weak, if at all, mechanical response. Peak temperatures at different frequencies have been included in the Arrhenius map of Figure 6. DMA traces correspond well to the DRS ones and, in agreement with them, do not show significant changes in time scale upon addition of POSS, with the exception of 10 wt % hybrid which shows slightly slower dynamics and, strikingly, 8 wt % which shows rather faster dynamics. DMA traces in Figure 6 appear at frequencies almost an order of magnitude higher than the extrapolated trajectory of the DRS traces. This is due to the different moieties probed by the two techniques, different temperature protocol (isothermal vs linear heating measurements), and finally the fact that dielectric loss is a compliance as opposed to the loss modulus. A last observation on the thermograms in Figure 9 we would like to point out is that the E″ peak at 50 °C is significantly suppressed upon addition of more than 4 wt % of particles in a nonsmooth manner. We cannot correlate this with the enthalpy of the corresponding MDSC endotherms as the latter are too weak to obtain reliable values. However, we may relate the phenomenon to a qualitative change in the morphology, e.g., a percolation of the POSS-extended phase. On the α′ Relaxation. Relaxations of the time scale of α′ accompanying the main α have been reported in polyurethane systems by us20,23,50 and in polyurea ones by Fragiadakis et al.51 The latter authors attributed α′ to unreacted amine groups acting as dielectric probes of the slowed down segmental mobility at interfaces, based mainly on the dependence of its strength on stoichiometry and the fact that it has no calorimetric or mechanical signature. In the present paper a weak contribution on glass transition signature on DSC thermograms and even weaker in DMA ones are present, but definitely the dielectric signature of α′ is much more evident in agreement with Fragiadakis et al.’s explanation, differing only in that here hydroxyl groups act as probes instead of amine. The slow mobility at interfaces is rather weak to be detected by thermal or mechanical techniques, but as it is probed by the strongly polar hydroxyl, it shows a significant dielectric response. The explanation is also compatible with our observations in ref 23 where α′ is enhanced when particles are tethered by urea linkages formed by amine functionalized vertex groups in the sense that amine groups are more polar than hydroxyl. Moreover, α′ is present also in the neat PTMEG but only

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CONCLUSIONS

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ASSOCIATED CONTENT

For the first time, a polyurethane system with POSS particles along its chain contour has been prepared and its molecular dynamics studied in detail. Small loadings (∼4 wt %) of POSS lead to two distinct phases. The first one is similar to the matrix with respect to chemistry and morphology. The second one is chain extended by the POSS and forms large but unstable crystalline structures. At high loadings (∼10 wt %) particles tend to form agglomerates; however, the aforementioned distinct phases are still visible in SEM images. Four techniques in agreement with each other show that time scale of segmental dynamics is not affected by the particles placed along the main chain. However, dielectric techniques and, to a lesser extent, calorimetry show that a slower component α′ is clearly enhanced and broadened in the presence of POSS. This is opposed to the increase of Tg without significant effect on α′ caused by tethering of particles as side groups. The significance of network topology on the properties of hybrids is evidenced, especially in complex architectures with intrinsic morphological features like polyurethanes and their microphase separation. We provide evidence that α′ reflects the dynamics of macrodiol chains close to heavy and immobile structures like particles or large and rigid hard microdomains. These dynamics are the same regardless of the nature of the heavy structure. Rubbery modulus is significantly reduced in the hybrids showing the poor mechanical properties of the POSS extended phase. It will be interesting in future work to examine the effects of POSS in different binding modeschain topologies like octafunctional POSS acting as heavy chemical cross-links.

S Supporting Information *

Figure S1 (backscattered electron image of a “stray” POSS crystal located in the bottom of a pore and its respective EDS spectrum); Figure S2 (dielectric loss spectra recorded with the 8 wt % hybrid); Figure S3 (temperature dependence of the shape exponent aCC for the matrix and two selected hybrids). This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

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

The authors declare no competing financial interest. H

dx.doi.org/10.1021/ma401417t | Macromolecules XXXX, XXX, XXX−XXX

Macromolecules

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ACKNOWLEDGMENTS K.N.R. and K.P. acknowledge partial funding by the National Science Centre in Poland under Contract No. DEC-2011/02/ A/ST8/00409. D.A. and P.P. acknowledge partial support by the European Union (European Social Fund - ESF) and Greek national funds through the Operational Program “Education and Lifelong Learning” - Research Program Aristeia. Authors are indebted to Ms Joanna Pagacz (CUT) for assistance with DMA measurements and to Dr. Daniel Fragiadakis for making and providing the software Grafity (grafitylabs.com), which we used for the analysis of dielectric spectra.

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dx.doi.org/10.1021/ma401417t | Macromolecules XXXX, XXX, XXX−XXX